Andrzej Rybak
email: andrzej.rybak@pl.abb.com

Karolina Gaska

Czeslaw Kapusta

Francois Toche

Vincent Salles

François Toche

Epoxy composites with ceramic core-shell fillers for thermal management in electrical devices

Keywords: functional composite, core-shell filler, thermal management, thermal conductivity

Introduction

Filled epoxy resins are well known and extensively used as insulating materials in power product applications. The addition of filler particles into epoxy resin can result in a significant improvement in the thermal, electrical or mechanical properties of the obtained composite [1][2][3][4][5]. The management of heat dissipation is a crucial consideration when ensuring the proper operation of such devices, with improper heat management potentially leading to malfunctions [6][7][8]. Optimized thermal management in electronic components, can lead either to reducing of the footprint or allowing a component to be rated for greater power ranges.

Besides the matrix thermal conductivity, the thermal properties of a composite are also related to the kind of filler used, the grain size, the content of this filler and its alignment in the matrix [1][2][3][4][9][10][11][12]. The quality of the filler-epoxy matrix interface seems to be crucial in ensuring the precise control of the thermal properties of a composite [13,14]. In particular, it is necessary to consider the strong phonon scattering processes as well as the interface resistance [2,15,16]. To increase thermal conductivity in epoxy composite materials, conductive paths must be maximized and thermal contact resistances must be reduced. A promising approach for creating composites that fulfil these requirements involves the modification of filler particles [17,18].

The heat dissipation properties of thermally conducting composites has been the subject of numerous theoretical and experimental studies [2][3][4]14]. Composites filled with high thermal conductivity materials have been shown to offer huge potential as a thermally conducting material. Whilst the composites offer significant improvements in effective thermal conductivity, there are also some disadvantages associated with using the material. First of all, fillers with very high thermal conductivity (HTC) are usually expensive, resulting in their application being questionable from an economical point of view, especially when high filler contents are required. On the other hand, less expensive fillers are not suitable for significant enhancement of thermal conductivity. A second disadvantage relates to the electrical insulation function of the composite. Requirements that a composite should maintain electrical insulation limits the variation in fillers that may be used, with ceramic particles with good dielectric properties being the most appropriate. Therefore, from an application point of view it is necessary to strike a compromise between the overall cost of composite and the type of filler necessary in order to obtain a suitable thermal conductivity for the whole system. A potential solution to these conflicting requirements is to create a core-shell like structure, whose core is composed of cheap standard material and covered with a low amount of expensive HTC shell [START_REF] Michalik | An insulating composite material with a polymer matrix[END_REF]. Several examples of different core-shell particles have been previously described [20][21][22][23]. Recently, the authors have demonstrated a successful approach with a fully ceramic core-shell fillers synthesis by means of carbothermal reduction and nitridation process which were incorporated into epoxy matrix in order to create electrical insulation with enhanced thermal conductivity [24].

In this work a polymer derived ceramics (PDCs) technique has been applied in order to produce a filler whose core is composed of silica flour, covered by boron nitride and silicon nitride shell. The PDCs is a known synthesis method for creating advanced ceramic materials [25][26][27][28][29].

The aim of the present work is to use a core-shell filler structure, with the core made of a standard and cheap material such as SiO 2 , and the coating made of a high-thermalconductivity (HTC) material, namely boron nitride (BN) or silicon nitride (Si 3 N 4 ) in order to increase the thermal conductivity of the filled epoxy resin systems. The addition of the modified filler can significantly increase the effective thermal conductivity of composite. The core-shell filler materials obtained were investigated using infrared (IR), X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), and also characterized by scanning electron microscopy (SEM) coupled with energy dispersive X-ray spectroscopy (EDS). A representative set of epoxy matrix samples filled with core-shell material has been prepared and compared with the reference samples. Thermal conductivity measurements have been performed on these samples at room temperature and also as a function of temperature over a wide range from 10K to 400K.

Experimental

Core-shell filler preparation

The first type of filler was prepared to form silicon nitride-coated silica, referred hereafter as SiO 2 @Si 3 N 4 . A volume of 40 mL of Perhydropolysilazane (PHPS, Sigma Aldrich), used as a silicon source, was poured into a beaker under ambient conditions. A folded filter containing 0.66 g of silica flour (Millisil W12, delivered by Quarzwerke) was consecutively dived into the solution during 15 minutes. The filter was then removed from the solution and unfolded in order to allow the powder to dry. The coated powder was left inside the unfolded filter at room temperature overnight under a hood.

After drying, the Si 3 N 4 -coated powders were transferred to an alumina crucible before a thermal treatment was applied at 1000°C with a rate of 5°C/min and a plateau of 2h under a flowing nitrogen atmosphere. A part of this powder was then heat-treated at 1400°C in order to obtain a crystallized nitride phase. A simplified scheme of the synthesis process is shown in Fig. 1

(upper part).

A second kind of filler was prepared to form a boron nitride coated silica (referred as SiO 2 @BN). Silica particles were coated with a layer of boron nitride (BN) starting from boric acid (H 3 BO 3 ) as a boron source. In a typical experiment, 200 mL of a 1M solution of boric acid in methanol was prepared at room temperature, before the addition of 0.61g of silica flour powders. The mixture was then stirred for 10 minutes before being kept for 5 days at room temperature under a hood until the methanol had totally evaporated. The dry powder was then heated in an alumina crucible to 1000°C at a rate of 10°C/min and held for 2h under a flowing ammonia atmosphere. The flow was finally replaced by a nitrogen one during the cooling step (cf. middle part of scheme shown in Fig. 1).

Composites preparation and thermal conductivity measurement technique.

An epoxy derived from bisphenol A (DGEBA) CY 228-1 which had been hardened with HY918 (derived by Huntsman) was selected as a standard system used for impregnation and casting of the power products. Composite samples were prepared by mixing the epoxy resin system with a filler content of 31 vol.%. Such a volume fraction was selected as typically used in the electrical insulation systems. A conventional vacuum casting method utilizing a typical curing profile found in the producer data sheet was used during the preparation of the samples (see bottom part of Fig. 1).

The thermal conductivity of the material was investigated using the pulse-power method of the thermal transport option of the Physical Property Measurement System (PPMS) equipment (Quantum Design). Small cylindrical samples with a diameter of 6 mm and a height in the range of 3 to 6 mm were cut from the material and tested. Measurements were carried out at room temperature for three samples of each composite, and the average values are presented. Heat pulses were injected into the test samples by a heater which was connected to a copper electrode that had been adhered to the sample surface. The thermal conductivity was measured under high vacuum (10 -6 Torr) in a steady-state at room temperature, but also as a function of temperature, starting from cryogenic temperatures, specifically 2K [4,30].

Results and discussion

Characterization of powders

Infrared spectroscopy

The obtained BN and PHPS coated SiO 2 powder was investigated by Fourier Transform Infrared spectroscopy (Attenuated Total Reflectance FTIR Nicolet 380, Thermo Scientific), in order to find the bonding character of boron and nitrogen with the silica particles.

The mid-infrared (MIR) spectra of pure and BN coated SiO 2 are shown in Fig. 2 (bottom and middle spectra). The positions and intensities of predominant bands in both cases are almost the same and are related to stretching vibration of Si-O (950-1200 cm -1 ) and silicon oxide bending modes (below 600 cm -1 ) [31][32][33][34]. However, a clear difference between the these two spectra is observed at 1380 cm -1 (broad band), where BN absorption is expected. Indeed, pure BN has two characteristic bands: a strong band at 1373 cm -1 attributed to B-N stretching vibrations and a weaker band at 816 cm -1 attributed to B-N-B bending vibrations [34].

In the case of MIR spectra for PHPS-coated silica before heat treatment (c.f. top spectrum in Fig. 2), the positions and intensities of SiO 2 bands for pure and coated silica are almost the same, but a difference shift is observed in the range 850-1000 cm -1 , where the strongest Si-N absorption bands are expected, with a Si-N-Si asymmetric stretching given at different values and ranges depending on the authors (925 cm -1 [35], 800-1000 cm -1 [36], 840-1020 cm -1 [37] .

This shift can indicate that the sample contains PHPS before the furnace treatment.

Electron microscopy imaging

A scanning electron microscope (SEM, Hitachi S800) equipped with energy-dispersive spectroscopy analysis (EDS) was used to analyze the coated particles. Comparing the EDS measurements of the pure SiO 2 powders with equivalent measurements of the material obtained after BN deposition, (see Fig. 3) clearly shows the presence of BN in the batch, as evidenced by the two peaks around 0.2 and 0.4 keV.

The SEM picture of SiO2@BN (see middle image in Fig. 4) seems to indicate agglomerates which probably appear due to a high content of BN since the B/Si molar ratio was fixed at 20 in our process.

The right image in Fig. 4 presents a SEM picture of SiO 2 @Si 3 N 4 powders after treatment at 1000°C. One can distinguish a film on the powder surface which could correspond to the Si 3 N 4 deposition already detected using FTIR technique. The particle size measured here is similar to that of the raw powder (see left image in Fig. 4) confirming there is no significant modification of the particle size after the formation of nitride layer.

X-ray photoelectron spectroscopy (XPS)

X-ray photoelectron spectroscopy (XPS) experiments were carried out on both batches: SiO 2 @BN and SiO 2 @Si 3 N 4 . In both cases a nitrogen element was detected, as well as boron for the boron containing powder. The obtained values after quantification are presented in Table 1, with a higher content of nitrogen for the BN-containing batch, meaning that the BN film seems to be thicker than the Si 3 N 4 one. The presence of carbon is due to the carbon tape on which the powders where deposited for SEM observations. These measurements were performed after grinding of the as-pyrolysed powders, implying that a part of the silicon detected could correspond to the silica. Moreover the coating is certainly not thick enough to prevent the electron beam from reaching the silica grains. This explains why, for instance, a high content of Si was detected in the BN-coated material.

X-ray diffraction

The core-shell powders were investigated by X-ray diffraction (PANalytical MPD). Boron nitride remains amorphous at temperatures lower than 1400°C. Therefore a trial was conducted where the temperature of a part of the SiO 2 @BN sample was increased up to 1400°C. However the XRD results highlighted only the formation of Crystoballite instead of a Quartz phase for SiO 2 . The main BN crystalline phase peak (around 2θ = 26.7°) was still absent.

The XRD patterns of as-received SiO 2 and treated SiO 2 are compared in Fig. 5. The crystalline phase of SiO 2 was pure Quartz. This phase was also observed in the sample after treatment 1400°C, but a second phase of SiO 2 (Cristobalite) was formed during the treatment.

The diffraction peak at positions 13° 2 and shoulder (at 31° 2) could be attributed to Si 3 N 4 . From the XRD patterns obtained after intermediate treatment at 1000°C, it may be observed that that the signal relating to Cristobalite is absent. This is consistent with the fact that silicon nitride crystallizes more rapidly above 1200°C. It must be noted that the hightemperature treatment was performed in order to highlight the formation of crystalline Si 3 N 4 but, obviously, this sample was subjected to a pyrolysis up to 1000°C in order to prevent the Quartz-to-Cristobalite transformation.

EDS analysis

The obtained powders were characterized by means of SEM coupled with EDS analysis. The modified core-shell grains have been dispersed in the epoxy resin and were polished with sub-micron diamond pastes what allowed the cross-section of each sample to be observed.

Several grains were analyzed precisely in order to check reproducibility of the layer formation. Fig. 6 shows an example of SEM image of a grain cross-section of the SiO 2 @BN sample.

Elemental analysis was conducted for each of the selected points as indicated in Fig. 6: points from 1 to 4 are placed on the edge of the grain, whereas point 5 is located in the center of the grain. Table 2 shows detailed results of EDS point analysis. It is clearly seen that content of boron and nitrogen is significantly increased at the outer layer of the grain which correlates well with the XPS results and confirms that the BN layer was synthesized successfully. In order to estimate the thickness of the BN coating, EDS linear analysis was performed on several grains. An example of such analysis is shown in Fig. 7.

After analysis of several EDS spectra it can be concluded that the layer of BN created on SiO 2 grain is not uniform, however the thickness is in the following range 0.8-3.0 m. Such a thickness should be sufficient to increase the effective thermal conductivity of the core-shell grain.

In the case of the SiO 2 grains coated with Si 3 N 4 it was hard to detect the coating. This observation is in good agreement with XPS data where the signal corresponding to Si 3 N 4 was much weaker than the equivalent for BN. Therefore one can conclude that layer of Si 3 N 4 created on SiO 2 is extremely thin.

Thermal conductivity results for composites

Thermal conductivity values measured for the prepared composites with incorporated coreshell fillers are presented in Table 3. The effective thermal conductivity was enhanced by 8% in case of SiO 2 @BN, whereas in the case of SiO 2 @Si 3 N 4 the observed thermal conductivity was significantly increased by 27% in relation to the reference sample. It is worth noting that although a much thinner layer of HTC coating is present, namely silicon nitride in the case of SiO 2 @Si 3 N 4 powder, the observed thermal conductivity is more enhanced (increased by three times) in comparison to SiO 2 @BN filler for which a thicker HTC coating, namely boron nitride, was obtained. The XRD results showed that crystallinity is poor for both the BN and Si 3 N 4 coatings, however a slightly better crystalline degree was observed for the silicon nitride shell. Therefore a higher enhancement of thermal conductivity can be expected in the case of SiO 2 @Si 3 N 4 due to the hindered scattering of phonons during heat transport. The observed results indicate that there is still place for optimization of the proposed core-shell filler structure, e.g. control of HTC layer thickness, its crystallinity and homogeneity.

Taking into consideration the results reported in literature for core-shell filler based polymer composites with similar filler content (40 wt.% for literature results vs. 50 wt.% in our case), e.g. alumina-coated graphene sheet (GS@Al 2 O 3 ) or reduced graphene oxide-encapsulated SiO 2 (SiO 2 @RGO) [21,22], the obtained values (0.57 and 0.45 W/mK, respectively) are comparable or even lower to those which were obtained in our study (see Table 3). On the other hand, one can find example of composites with incorporated core-shell structure giving better results, e.g. silica coated silver (Ag@SiO 2 ) nanoparticles [20], however the fraction of filler was much higher than in our case (50 vs. 31 vol.%), what is an aspect which could have a strong influence on the results. It might be also hard to justify the use of noble metal nanostructures for industrial applications due to problems associated with volume production and the handling of nanoparticles.

An additional experiment, namely the temperature dependence of the thermal conductivity, was also conducted, with the results being presented in Fig. 8. One can observe that the effective thermal conductivity of composites is stable over a wide range of temperatures from 300K to 150K. This makes the composites filled with the modified core-shell filler very promising from the application point of view, because stable thermal conductivity is necessary for the correct operation of electrical devices.

The schematic representation of the heat transport mechanism is shown in the right part of Fig. 8. The HTC shell material is mainly responsible for effective heat dissipation and enhancement of the overall thermal conductivity of the composite.

Conclusions

The performed research work has demonstrated that a new type of core-shell fillers can be prepared by means of the polymer derived ceramics synthesis method. On the basis of the obtained characterization results for the modified particles it was confirmed that the coreshell structure formation was successful. It was shown that the application of the core-shell fillers, especially silica with silicon nitride shell (SiO 2 @Si 3 N 4 ), has resulted in an improvement of the effective thermal conductivity of the filled epoxy composite. The enhanced thermal conductivity (up to 27%) of the system with incorporated novel core-shell filler can open the way for more affordable epoxy composites suitable for application as an advanced material for the thermal management in electrical devices.

Tables

Figure captions:Figure 1 .

 1 Figure captions:

Figure 2 .

 2 Figure 2. MIR spectra of reference SiO 2 , BN coated SiO 2 (SiO 2 @BN), and PHPS-coated

Figure 3 .

 3 Figure 3. Results of the EDS measurements for pure SiO 2 and SiO2@BN powders.

Figure 4 .

 4 Figure 4. SEM image of raw SiO 2 (left), SiO 2 @BN (middle) and SiO 2 @Si 3 N 4 (right)

Figure 5 .

 5 Figure 5. XRD patterns of SiO 2 as-received (top graph) and modified SiO 2 @Si 3 N 4 particles

Figure 6 .

 6 Figure 6. SEM image of cross-section of SiO 2 @BN grain. White numbers indicate the

Figure 7 .

 7 Figure 7. SEM image of cross-section of SiO 2 @BN grains (top image). The white line

Figure 8 .

 8 Figure 8. Temperature dependences of thermal conductivity for composites with core-shell

Figure 1 .

 1 Figure 1.

Figure 4 .

 4 Figure 4.

Figure 8 .

 8 Figure 8.

Table 1 .

 1 : Composition determined during XPS analysis.

	Content (wt.%)	B	N	C	O	Si
	SiO 2 @BN	34,3	33,9	13,4	15,2	3,3
	SiO 2 @Si 3 N 4	0,0	5,2	9,7	56,3	28,8

Table 2 .

 2 Results of EDS point analysis of cross-section of SiO 2 @BN grain. Points correspond to places indicated in Fig.7.

	Point	B (wt. %)	N (wt.%)
	1	21.40	8.69
	2	20.40	5.80
	3	17.71	13.06
	4	20.11	11.79
	5	8.39	1.57

Table 3 .

 3 Thermal conductivity results obtained at 300K, for composites filled (31 vol.%) with designed fillers SiO 2 @BN and SiO 2 @Si 3 N 4.

	Sample	(W/mK)	Relative increase
		at 300K	(%)
	SiO 2	0.62	-
	SiO 2 @BN	0.67	8
	SiO 2 @Si 3 N 4	0.79	27

Acknowledgements

The authors gratefully acknowledge the support from James Ottewill for correcting English language.


Adrien Leroy

Sofia Ribeiro

Carole Grossiord

Antoine Alves

Robert H Vestberg

Vincent Salles

Céline Brunon

Kerstin Gritsch

Brigitte Grosgogeat

Yves Bayon
email: yves.bayon@medtronic.com

FTIR microscopy contribution for comprehension of degradation mechanisms in PLA-based implantable medical devices

Keywords:

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

Introduction

Polylactic acid (PLA) is an aliphatic polyester that can be synthesized from lactic acid or by ring-opening polymerization of lactide. It has two optical isomers, L-and D-isomers, the first being the natural occurring polymer. Poly(Llactic) acid (PLLA) is a semi-crystalline polymer with a glass transition temperature of 60-65 °C and a melting temperature approximately 175 °C [START_REF] Castillo | Crystallization and morphology of biodegradable or biostable single and double crystalline block copolymers[END_REF][START_REF] Nair | Biodegradable polymers as biomaterials[END_REF][START_REF] Garlotta | A literature review of poly (lactic acid)[END_REF]. PLLA can have different crystal structures. The most common structure is α form which is produced by polymer melt or solution crystallization. Another form is the β form which comes from crystallization at high temperature under high draw ratios as for example in fiber spinning [START_REF] Saeidlou | Poly (lactic acid) crystallization[END_REF].

PLLA has attracted huge interest from the biomedical and pharmaceutical fields and has been massively used as a material in surgical operations, tissue regeneration and drug delivery systems, for decades now [START_REF] Ikada | Biodegradable polyesters for medical and ecological applications[END_REF]. Besides being considered as a "green polymer", PLLA's great advantage is that it is degraded, mostly by hydrolysis [START_REF] Zhou | Nanoclay and crystallinity effects on the hydrolytic degradation of polylactides[END_REF][START_REF] Krikorian | Crystallization behavior of poly (L-lactic acid) nanocomposites: nucleation and growth probed by infrared spectroscopy[END_REF]. In tissue engineering applications it is vital to fully understand the degradation process of the biomaterial in question, since the coordination of molar and gravimetric mass losses and tissue healing determines how successful the tissue repair is. The hydrolytic degradation of PLA occurs by bulk erosion process where diffusion-reaction mechanisms are in play [START_REF] Grizzi | Hydrolytic degradation of devices based on poly (DL-lactic acid) size-dependence[END_REF]. The human body metabolizes this polymer by hydrolysis into lactic acid, which is further broken down to water and carbon dioxide and eventually excreted by the organism [START_REF] Castillo | Crystallization and morphology of biodegradable or biostable single and double crystalline block copolymers[END_REF][START_REF] Nair | Biodegradable polymers as biomaterials[END_REF]. This type of polymer has a degradation rate of 50% in about 1 to 2 years [START_REF] Ikada | Biodegradable polyesters for medical and ecological applications[END_REF].

It has been previously reported many studies about PLA degradation in vitro, reporting the monitoring of parameters such as crystallinity [START_REF] Gorrasi | Effect of PLA grades and morphologies on hydrolytic degradation at composting temperature: assessment of structural modification and kinetic parameters[END_REF][START_REF] Tsuji | Properties and morphology of poly (L-lactide) 4. Effects of structural parameters on long-term hydrolysis of poly (L-lactide) in phosphate-buffered solution[END_REF][START_REF] Renouf-Glauser | hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 25/27 of PLLA[END_REF], temperature [START_REF] Gorrasi | Effect of PLA grades and morphologies on hydrolytic degradation at composting temperature: assessment of structural modification and kinetic parameters[END_REF][START_REF] Xu | Effects of temperature and pH on the degradation of poly (lactic acid) brushes[END_REF][START_REF] Li | Hydrolytic degradation characteristics of aliphatic polyesters derived from lactic and glycolic acids[END_REF][START_REF] Oliveira | The role of shear and stabilizer on PLA degradation[END_REF], time [START_REF] Grizzi | Hydrolytic degradation of devices based on poly (DL-lactic acid) size-dependence[END_REF][START_REF] Li | Effect of reinforcing particles on hydrolytic degradation behavior of poly (lactic acid) composites[END_REF][START_REF] Partini | Determination of crystallinity of an aliphatic polyester by FTIR spectroscopy[END_REF][START_REF] Tsuji | Synthesis and hydrolytic degradation of substituted poly (DL-lactic acid)s[END_REF][START_REF] Tsuji | Autocatalytic hydrolysis of amorphous-made polylactides: effects of L-lactide content, tacticity, and enantiomeric polymer blending[END_REF] and pH [START_REF] Xu | Effects of temperature and pH on the degradation of poly (lactic acid) brushes[END_REF][START_REF] Li | Further investigations on the hydrolytic degradation of poly (DL-lactide)[END_REF][START_REF] Lostocco | The hydrolysis of poly (lactic acid)/poly (hexamethylene succinate) blends[END_REF][START_REF] Maitz | Applications of synthetic polymers in clinical medicine[END_REF][START_REF] Schliecker | Characterization of a homologous series of D, L-lactic acid oligomers; a mechanistic study on the degradation kinetics in vitro[END_REF]. The degradation products of PLA materials cause a drop of pH, contributing to a faster degradation and, when implanted, a locally inflammatory response persisting until the full disappearance of the materials [START_REF] Cheung | A critical review on polymer-based bio-engineered materials for scaffold development[END_REF].

Most studies on PLLA hydrolysis are performed at 37 °C (standard temperature for in vitro assays). However, some studies present the time-temperature equivalent relationship degradation process of PLLA, at temperatures up to 90 °C [START_REF] Santis | Nucleation and crystallization kinetics of poly (lactic acid)[END_REF]. Since PLLA takes a long time to completely degrade in real aging conditions (i.e. at 37 °C), accelerated degradation can generate other PLLA materials, in about a week, which are good representatives of the aging process, occurring at physiological temperature.

Crystallization in biodegradable polymers is a vital process due to its influence over physical, mechanical and degradability properties [START_REF] Santis | Nucleation and crystallization kinetics of poly (lactic acid)[END_REF]. There are many studies on the crystallization kinetics of PLLA using techniques such as DSC, NMR and FTIR [START_REF] Gorrasi | Effect of PLA grades and morphologies on hydrolytic degradation at composting temperature: assessment of structural modification and kinetic parameters[END_REF][START_REF] Oliveira | The role of shear and stabilizer on PLA degradation[END_REF][START_REF] Partini | Determination of crystallinity of an aliphatic polyester by FTIR spectroscopy[END_REF][START_REF] Santis | Nucleation and crystallization kinetics of poly (lactic acid)[END_REF][START_REF] Migliaresi | The effect of thermal history on the crystallinity of different molecular weight PLLA biodegradable polymers[END_REF][START_REF] Pan | Temperaturevariable FTIR and solid-state 13C NMR investigations on crystalline structure and molecular dynamics of polymorphic poly (L-lactide) and poly (L-lactide)/poly (D-lactide) stereocomplex[END_REF]. Compared to other techniques, FTIR spectroscopy shows advantages in the analysis of the polymer crystallinity. FTIR is a non-destructive fast technique, hence suitable for a rapid determination of crystallinity and crystal structure. However this technique presents also drawbacks such as making the correct identification of bands to the amorphous and crystalline content, and find the accurate absorption coefficient ratio between amorphous and crystal phases [START_REF] Partini | Determination of crystallinity of an aliphatic polyester by FTIR spectroscopy[END_REF].

FTIR microscopy should also present key advantages in confirming the presence or absence of residual materials by their chemical nature and IR absorption profile. It should be particularly useful at advanced stages of degradation, associated with the fragmentation and possible migration of the implanted materials. The aim of this study is to develop a new imaging technique to identify the implanted material and to track the degradation of resorbable polymers in vivo. The crystallinity of the material was observed using FTIR microscopy analysis, giving the characteristic spectra of PLLA at different degradation levels after aging at 90 °C. The in vitro degradation profile was analyzed and then the imaging technique was then translated to in vivo studies http://eproofing.springer.com/journals/printpage.php?token=VY4BC44hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 5/27

with analysis on fixed tissue explant sections. Finally, the potential of this FTIR imaging technique has been applied to the degradation of other resorbable polyesters.

Materials and methods

Materials

Tetrahydrofuran (99.9 %) (THF) was purchased from Sigma Aldrich and Phosphate-Buffered Saline (PBS, pH 7.4) from ThermoFisher Scientific (Gibco).

Two kinds of bioresorbable hernia meshes were used for the degradation studies.

Poly(L-lactide) (PLLA) meshes (Sofradim Production, France) were used for both in vitro and in vivo degradation studies and TIGR ® Matrix Surgical Meshes (Novus Scientific, Sweden) (thereafter called TIGR mesh) were only used for the in vivo study. TIGR mesh is knitted from two different synthetic resorbable fibers: the first kind of fibers is made of a copolymer of lactide (LA), glycolide (GA) and trimethylene carbonate (TMC) and the second one is made of a copolymer of LA and TMC.

Characterizations

Molecular weight

The number average molecular weight (M n ), mass average molecular weight (M w ) and dispersity (Đ) of the polymers were determined by Size Exclusion Chromatography (SEC) using a Waters Aquity UPLC fitted with two Shodex GPC KF-805L columns, a Shodex GPC KF-G guard column and Wyatt refractive index and multiangle light scattering detectors. The mobile phase was THF, flowing at 1 mL min -1 . For each tested polymer, three samples were analyzed with two injections (3.5 mg/mL).

Differential Scanning Calorimetry (DSC)

DSC measurements were carried out on a Mettler Toledo DSC thermal analyzer.

For each reference, four samples (5-6 mg) were submitted to a heating scan to 200 °C (10 °C/min), a cooling scan to 0 °C (10 °C/min) and a second heating scan to 200 °C (10 °C/min). Glass transition temperature (T g ), melting temperature (T m ) and degree of crystallinity (X c ) were determined from the second heating ramp. A reference enthalpy of melting of 93.1 J/g was used to calculate the crystallinity of PLA [START_REF] Fischer | Investigation of the structure of solution grown crystals of lactide copolymers by means of chemical reactions[END_REF]. FTIR analyses were performed, using a Perkin-Elmer Spotlight 400 microspectrometer, on polymer samples obtained from the in vitro study and on the tissue sections obtained from in vivo experiments. FTIR mapping was performed in Attenuated Total Refection (ATR) mode using the ATR-imaging accessory with a germanium crystal. The scanned area was 400 × 400 µm² (lateral resolution of 1.56 µm). The MCT (Mercury-Cadmium-Telluride) detector covered the range from 720 to 4000 cm -1 with a spectral resolution of

In vitro degradation study

For this degradation study, polymer samples (15 × 5 cm² PLLA meshes) were initially weighed and placed in PBS (two meshes in 100 mL) at 90 °C. The PBS was changed every one or 2 weeks. At each scheduled time point, samples were removed from PBS and rinsed 3 times with distilled water. After at least 12 h of drying under vacuum, they were weighed to calculate the gravimetric mass loss, analyzed by SEC and DSC, and FTIR spectra were finally obtained.

Preparation of pre-degraded samples for in vivo degradation study

Both PLLA and TIGR meshes undergo slow degradation by bulk hydrolysis and are expected to be completely absorbed 36 months after implantation [START_REF] Middleton | Synthetic biodegradable polymers as orthopedic devices[END_REF]. The assessment of the entire in vivo degradation process of such devices therefore requires a long follow-up. To decrease this follow-up duration, samples were submitted to an in vitro pre-degradation process before implantation. Based on the data obtained from in vitro degradation studies, some samples (2 × 1.5 cm²) were pre-degraded in PBS at 50 °C so that we have samples in three different degradation states : non pre-degraded, partially pre-degraded and highly predegraded. The properties of these samples before implantation are listed in Table 1. All samples were provided as sterile meshes ready to implant (2 × 1.5

http://eproofing.springer.com/journals/printpage.php?token=VY4BC44hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 7/27 cm²) except highly pre-degraded PLLA samples which were provided as sterile moistened particulates (15-20 mg) corresponding to residues of meshes (2 × 1.5 cm²) recovered after a few seconds of centrifugation at 1000 rpm. 1).

Surgical procedure

Buprenorphine (Buprecare ® , Axience, 0.02 mg/kg) was administered subcutaneously prior to surgery. Rabbits were anesthetized by intramuscular injection of a mixture of ketamine hydrochloride (Ketamine 1000, Virbac, 20 mg/kg) and xylazine (Rompun ® , Bayer, 5 mg/kg). When requested, anesthesia was maintained by inhalation of an O 2 -isoflurane mixture (IsoFlo ® , Axience, 0.5-5%). After the anesthetic has taken effect, sufficiently large areas of the rabbits were shaved using an electric clipper with a fine comb, avoiding any cutaneous trauma. The shaved area was cleaned with povidone iodine (Vetedine ® Savon, Vetoquinol), wiped with 70% isopropyl alcohol, painted with povidone iodine solution (Vetedine ® Solution, Vetoquinol) and draped. Eight incisions (four per side) large enough to accommodate each sample were made through the skin and parallel to the vertebral column (approximately 1.5 cm from the midline, 1.5-2 cm length). Pockets were formed by blunt dissection in the subcutaneous tissue with a standardized depth (approximately 3.5 cm). Before implantation, two non-absorbable stitches (Surgipro TM II, 4-0, Covidien) were placed approximately 8 mm above the implanted article and approximately 8 mm below the cutaneous incision, in the subcutaneous tissue towards the skin, spaced from approximately 1.5 cm to locate the articles at termination. Four PLLA or TIGR mesh articles were introduced in the pockets of one side with one article per pocket, articles were placed at appropriately spaced intervals (approximately 2 cm). Highly pre-degraded PLLA samples (particulates) were placed in their subcutaneous pockets using a spatula. On the opposite side, four test and control articles were similarly implanted. The skin was closed with stainless steel skin staples and pressure bandages were applied. The rabbits were monitored until recovery from the anesthetic procedure and then returned to their respective cages.

Termination and histologic preparation

Four, 26 or 52 weeks after implantation, the designated rabbits were euthanized by intravenous injection of pentobarbital (Dolethal ND , Vetoquinol). The subcutaneous tissues were excised, sampled and flat fixed in 10% neutralbuffered formalin (NBF). After complete fixation, the tissue explants were dehydrated in alcohol solutions of increasing concentration, cleared in xylene and embedded in paraffin. For each sample, two central longitudinal ventrodorsal cross sections were prepared using a microtome (MICROM ® , France).

One was used for histological analysis (4-7 µm thick) and one (9 µm thick) was used for FTIR mapping.

Histology

Sections were stained by Safranin-Hematoxylin-Eosin and modified Masson's Trichrome.

FTIR microscopy

Three different FTIR maps of tissue explant sections (≈10 × 20 mm²) after dewaxing, were processed. Components were identified and colored thanks to Principal Components Analysis (PCA). R2 ratio (1130/1080 cm -1 band intensity) and amide I protein (1646 cm -1 ) were imaged in false colors (low values appear in blue, high values appear in red).

At explantation, parts of the samples where used for SEC analyses. The explanted sample was exposed to a collagenase solution (1 mg/ml) in PBS (pH = 7.4) for 16 h. The samples where then rinsed three times with water and dried under vacuum for 12 h before SEC analysis.

Results

In vitro degradation of PLLA meshes

The hydrolytic degradation process of aliphatic polyesters such as PLA or PGA occurs by random scissions of ester bonds within the polymer chains. Since this process may meet slow kinetics, especially for highly crystalline polymers like PLLA, an accelerated degradation protocol was used to study the behavior of our polymers. Indeed, it was calculated that the degradation process would be 496 times faster at 90 °C than at 37 °C [START_REF] Lyu | Kinetics and timetemperature equivalence of polymer degradation[END_REF]. At this temperature, the estimated degradation time of our PLLA meshes would be about 1 week (assuming full degradation within 4-5 years). The in vivo degradation is anticipated to be faster, notably at the latest stages, with more susceptibility to be degraded by the polymorphonuclear leukocyte cells by their chemical armors and endocytosis/phagocytosis of fragments [START_REF] Urayama | Microstructure and thermal properties of polylactides with different L-and D-unit sequences: importance of the Helical nature of the L-sequenced segments[END_REF].

To measure the evolution of FTIR peaks during PLLA degradation, spectra were obtained from non-degraded (0 h) and degraded [START_REF] Grizzi | Hydrolytic degradation of devices based on poly (DL-lactic acid) size-dependence[END_REF][START_REF] Santis | Nucleation and crystallization kinetics of poly (lactic acid)[END_REF] and 120 h at 90 °C) PLLA mesh samples ( Fig. 1 Evolution of FTIR spectra during PLLA degradation. Spectra were obtained from nondegraded (0 h) and degraded (8, 24 and 120 h at 90 °C) PLLA mesh samples. These spectra revealed notably peak shifts and an increase of intensity for bands at 1209 and 1130 cm -1 during degradation. See particularly Fig. 1c The IR spectra of all PLLA samples, non-degraded and degraded present the common features of PLA polymers formed of L-lactic acid units. In particular, a clear peak at 1208 cm -1 (νasCOC + CH 3 ), is observed whereas a faint peak is noticed around 1265 cm -1 (δCH + νCOC) which is typical of PLLA [START_REF] Urayama | Microstructure and thermal properties of polylactides with different L-and D-unit sequences: importance of the Helical nature of the L-sequenced segments[END_REF]. It is assumed that all analyzed PLLA samples in their native and degraded states are mostly in the α crystalline configuration. The IR spectra are similar to previously analyzed PLLA, notably with the lack the characteristic band of PLLA β at 908 cm -1 [START_REF] Zhang | Crystal modifications and thermal behavior of poly (L-lactic acid) revealed by infrared spectroscopy[END_REF]. The fact to confirm the homogeneity of the PLLA samples in terms of crystalline structure should facilitate the real contribution of the crystallinity changes occurring during the thermal hydrolytic degradation of the PLLA materials on their IR profile. Indeed, the corresponding spectra revealed peak shifts and an increase of intensity for bands at 1209 and 1130 cm -1 during degradation. These bands are related respectively to the asymmetric C-O-C stretching and asymmetric CH 3 rocking vibrations [START_REF] Kister | Effects of morphology, conformation and configuration on the IR and Raman spectra of various poly (lactic acid) s[END_REF]. It was already shown that such changes occurring with aliphatic polyesters could be explained by a change in the polymer crystallinity [START_REF] Partini | Determination of crystallinity of an aliphatic polyester by FTIR spectroscopy[END_REF][START_REF] Kister | Effects of morphology, conformation and configuration on the IR and Raman spectra of various poly (lactic acid) s[END_REF][START_REF] He | Novel FTIR method for determining the crystallinity of poly (ε-caprolactone)[END_REF]. In our case, one of the consequences of PLLA degradation is a decrease of its molecular weight caused by chains scissions. Shorter PLLA chains mobility is thus increased, which allows them to reach a more crystalline state, in the α configuration. Based on the literature [START_REF] Krikorian | Crystallization behavior of poly (L-lactic acid) nanocomposites: nucleation and growth probed by infrared spectroscopy[END_REF], we assumed that the changes observed for the bands at 1209 and 1130 cm -1 were caused by the increase of PLLA crystallinity and corresponding increase in the amount of α crystals during degradation. Hereby it would be possible to follow the degradation of the polymer by calculating peak intensity ratio differences. Ratios R1 (1209/1180 cm -1 band intensity) and R2 (1130/1080 cm -1 band intensity) were selected for this study.

To validate this hypothesis, we analyzed two non-degraded PLLA samples which were previously submitted to two different thermal treatments. The first sample was heated from 25 to 220 °C (10 °C/min) and slowly cooled to 25 °C (1 °C/min) to allow crystallization. The second one was heated from 25 to 220 °C (10 °C/min), removed from the oven and immediately immersed in liquid nitrogen to prevent crystallization and reach an amorphous state. After these treatments, both samples were analyzed by DSC: the first sample only exhibited a melting peak (ΔH = -41 J/g) whereas the second one exhibited a cold crystallization peak (33 J/g) followed by a melting peak (-32 J/g). These results confirmed that the first sample was able to crystallize during the slow cooling step whereas the second one was almost totally amorphous after being quenched. As a consequence, these two samples had different crystallinity rates but had similar chemical compositions.

For both these samples of crystallized and amorphous PLLA as well as for nondegraded (0 h) and highly degraded (120 h) PLLA mesh samples, ratios R1 and R2 were calculated (Table 3). Results showed that both ratios were higher with highly degraded PLLA than with non-degraded PLLA. According to our http://eproofing.springer.com/journals/printpage.php?token=VY4BC44hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 12/27

hypothesis, this would mean that ratios R1 and R2 increase with the crystallinity of the polymer. This was confirmed by the ratios calculated for crystallized PLLA (≈ non-degraded PLLA) and for amorphous PLLA (<non-degraded PLLA). 1 for the characteristics of the implanted meshes.) To verify that in vitro degradation was representative of the in vivo degradation SEC analyses of in vitro and in vivo degraded samples was performed. The results show that the degradation rate at +37 °C is similar (M w in vitro 26 weeks = 61 000 g/mol and M w in vivo 26 weeks = 57 000 g/mol). This confirms that in vitro degradation data can be representative for comparison with in vivo results, with however slightly faster in vivo degradation as anticipated as discussed in the previous section.

As shown in the in vitro degradation study, FTIR ratios R1 and R2 reflect the crystallinity and consequently the degradation of PLLA. To monitor the state and the integration of a medical device made of PLLA, an FTIR ATR-imaging study was realized on tissue sections using three different peak intensities or peak intensity ratios. First, the different components and especially polymer and tissue areas were identified using PCA mapping. Then, the degradation state of the polymer areas was monitored thanks to R1 or R2 ratios. Indeed, R1 ratio mapping showed less accuracy in case of highly degraded fibers surrounded with tissue because of partial interference between bands at 1208 cm -1 of PLLA and 1232 cm -1 of protein. Only R2 pictures were therefore analyzed in this study. Finally, a last mapping was realized at 1646 cm -1 (the amide I protein band) to detect the presence of neo-tissue extracellular matrix (ECM). Color scales for R2 ratio and 1646 cm -1 peak intensity are presented in Supplementary Fig. 2.

First, 0 PLLA samples implanted without in vitro pre-degradation were studied after 4 and 52 weeks implantation (Fig. 2). At the earliest stage of degradation, fibers did not seem damaged and were still easily visible by optical microscopy (OM). PCA mapping allowed even easier identification of the polymer within the tissue section and highlighted the fibers' structure conservation. On the other hand, R2 ratio mapping did not evidence important differences in the degradation state within the fibers. Slightly more crystalline areas could be seen along the outer edge of the fibers as well as in the middle of one of them but these differences could not be considered as major ones. After 4 weeks, degradation therefore seemed to have not visibly started. However, amide I protein mapping revealed early ECM deposition at the fibers interface.

Fig. 2 FTIR microscopy of non-pre-degraded PLLA meshes ( 0 PLLA) explants, at 4 and 52 weeks after subcutaneous implantation in rabbits. The intensities of R2 ratio, a degradation indicator based on increased crystallinity, and the amide I protein band is represented in false colors with a gradient from dark blue (low intensity) to hot red (high intensity). OM Optical microscopy, PCA Principal Component Analysis, R2 ratio: 1130/ 1080 cm -1 band intensity; Amide I protein band at 1646 cm -1

Fifty two weeks after implantation, it was still easy to locate PLLA fibers on both OM pictures and PCA maps and no material breaks or splits could be noticed. This time, however, a crystallinity gradient was observed on the R2 ratio map: a more crystalline (meaning more degraded) area appeared in the central part of each fiber. Even though these differences were still weak, they were in accordance with the current hypotheses on PLLA degradation mechanism. Indeed, PLLA is known to undergo bulk degradation by hydrolytic scissions of ester bonds in the polymer chains. This phenomenon is accelerated in the core of the material by the increasing presence of carboxylic acid-ended oligomers generated by the chain scissions which are not able to diffuse out of the material. As a consequence, these oligomers have an autocatalytic action and cause faster degradation in the inner part of the polymer device. Observations on the R2 ratio map after 52 weeks suggested that the PLLA fibers were at the beginning of this process. Besides, an ECM deposition could once again be observed at the interface of the fibers.

To simulate longer times of degradation, pre-degraded 1 PLLA samples were then studied 4, 26 and 52 weeks after implantation (Fig. 3). After 4 weeks, location of the fibers was still possible by OM but some areas seemed damaged and some parts were more difficult to distinguish from the surroundung tissues. PCA maps also revealed both undamaged and partially fragmented fibers, suggesting that the samples were actually in a more advanced degradation state than 0 PLLA samples 52 weeks after implantation. R2 ratio mapping also showed an increased overall crystallinity of this 1 PLLA sample after 4 weeks compared to 0 PLLA after 52 weeks. However, R2 ratio mapping showed unexpected results since the faster inner degradation could not be observed any more. Since the samples have been pre-degraded at 90 °C it is possible that the the degradation products can diffuse more easily when degrading at temperatures higher than the glass transition temperature (Tg = 50-60 °C). In addition, the acidification of the degradation medium over time could also mask the autocatalytic effect. After implantation, the autocatalysis phenomenon is probably too limited after such a short in vivo degradation time (autocatalysis effect was not observed for 0 PLLA samples after 4 weeks either). Structure damaging of the fibers could also distort FTIR mapping and generate unreliable results. As regards tissue colonization, amide I protein band mapping showed ECM deposition similar to that described for the previous samples at the interface of the fibers.

Fig. 3 FTIR microscopy of slightly pre-degraded PLLA meshes ( 1 PLLA) explants, at 4, 26 and 52 weeks after subcutaneous implantation, in rabbits. Legends are the same as for Fig. 2 http://eproofing.springer.com/journals/printpage.php?token=VY4BC44hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 15/27

After 26 weeks, OM pictures showed damaged samples with the appearance of cracks within the fibers, making them difficult to locate compared to all the previous samples. In this context, the use of PCA mapping proved to be very useful to precisely locate the polymer in the tissue section. It also clearly evidenced the presence of cracks in the material, which were particularly important since they were not visible on R2 ratio maps. Contrary to what was observed in 1 PLLA samples after 4 weeks, these maps revealed many highly degraded areas in the core of the fibers after 26 weeks. The overall degradation polymer in a more advanced state than after 4 weeks and the autocatalysis degradation areas is observed again. Interestingly, where cracks appeared on the PCA maps, R2 ratio mapping sometimes seemed to display low-degraded polymer areas. These differences confirmed that material damaging can cause misinterpretations and proved that R2 ratio mapping needs to be used and interpreted carefully and completed with other information to be reliable. Concerning the tissue growth, ECM deposition could once again be noticed around the fibers thanks to protein mapping. The ECM areas however seemed slightly wider than they were before, showing the integration of the device within the host tissues. Fifty two weeks after implantation, samples were so fragmented that it became very difficult or even impossible to recognize 1 PLLA residues in the tissue section by OM. As for 26 weeks explants, PCA mapping was necessary to locate the polymer and confirmed the high fragmentation of the material. R2 ratio mapping pictures proved that they should not be used alone for interpretation. The many cracks observed with PCA mapping were indeed not visible with R2 ratio mapping. Moreover, the material seemed to be homogeneously crystalline and less degraded than after 26 weeks, which is not coherent with the degradation mechanism of PLLA. All these inconsistencies could be explained thanks to the last FTIR maps. Indeed, protein mapping revealed an extensive presence of ECM all over the tissue section, including within the PLLA residues. This showed that, at this stage, the neo-tissue colonization may make the R2 ratio mapping provide wrong information and lead to misinterpretations concerning the degradation state of the polymer if it is not correlated with other FTIR mappings.

Finally, 2 PLLA explants were studied to evaluate the evolution of a highly predegraded sample (Fig. 4). On the one hand, OM and PCA maps observations 26 weeks after implantation were quite similar to those of 1 PLLA samples after the same in vivo degradation time. Indeed, the fibers were still visible but had suffered important damages. On the other hand, R2 ratio maps seemed to show that 2 PLLA was in a lower degradation state than 1 PLLA was at the same moment. Likewise it was observed for 1 PLLA after 52 weeks, protein mapping of 2 PLLA after 26 weeks revealed the presence of ECM around the fibers but also, to a lesser extent, within the material, which induced this incorrect information on R2 ratio maps. Once again, it was clear that R2 ratio maps had to be analyzed with caution and that information provided by the other FTIR mappings had to be taken in account for reliable interpretation. Tissue sections of 2 PLLA after 52 weeks are not presented in this study since the material was so degraded that yarn fragmentation, cell infiltration and ECM deposition made degradation monitoring impossible (situation close to that of 1 PLLA 52 weeks after implantation).

This first in vivo study proved that FTIR mapping was a useful and efficient technique to identify the implanted material and monitor the degradation and the resorption of PLLA medical device after implantation. It also exhibited the limits of this characterization method.

In vivo degradation of TIGR meshes

After demonstrating the feasability of locating PLLA fibers in tissue sections and monitoring their degradation thanks to FTIR mapping, we got interested in the possibility of applying the same procedure to other degradable polymers.

Attention was focused on TIGR meshes, whose fast resorbing (copolymer of GA with small amounts of LA and TMC) and slow resorbing (copolymer of LA with small amounts of TMC) multifilament fibers are both described as degrading by bulk hydrolysis [START_REF] Hjort | Three-year results from a preclinical implantation study of a long-term resorbable surgical mesh with time-dependent mechanical characteristics[END_REF]. As for the PLLA samples, some TIGR mesh were submitted to an in vitro pre-degradation step before implantation (Table 3).

First, a mapping test was performed on non pre-degraded samples ( 0 TIGR) after 4 weeks of in vivo degradation (Fig. 5). At this early degradation stage, OM pictures still revealed the presence of the two kinds of fibers (of two different sizes). By performing the same PCA treatment as for PLLA (PCA 1), it was possible to identify the different components of the section. The presence of LA units in slow resorbing fibers was still highlighted by a yellow to red coloration whereas the fast resorbing fibers were located by TMC spectrum correspondance and appeared in blue. Another PCA treatment (PCA 2) allowed even easier location of PGA-based fibers (colored in green) and brought to light a very early fracturation phenomenon within this material, suggesting that these fibers would indeed reach full resorption after a short time. Interestingly, it should be noticed that this PCA treatment also emphasized a third component (colored in blue) which could be identified as neo-tissue ECM thanks to amide I protein mapping. R1 and R2 ratios mapping were also performed on the tissue sections and seemed to allow degradation monitoring for slow resorbing fibers. On the contrary, PGA-based polymer was not visible on these maps since the peaks used for R1 and R2 ratios calculation were not present on this polymer's FTIR spectrum at this stage due to the high GA content (Supplementary Fig. 3), thus preventing any degradation monitoring for this material (at least with these ratios). In such case of multi-material degradable device implantation, PCA 2 treatment thus appeared to be a very efficient tool to locate the different components and to follow their resorption through time. The resolution of obtained images was even sufficient to distinguish individual fibers. However, R1 or R2 ratio mapping could be employed to follow crystallinity and degradation states only for the slow resorbing fibers made of PLA-PTMC copolymer. FTIR mappings of TIGR meshes at different degradation stages were then performed to evaluate the potential of this method to monitor the slow resorbing fibers degradation (Fig. 6). Four weeks after implantation, the PLA-based fibers of native 0 TIGR samples were still easily visible on both OM pictures and PCA (PCA 1) maps. However, PCA maps exhibited not only fibers perpendicular to the section plan but also some that seemed to be parallel to this plan and were thus hardly identified on the OM pictures. It can be noticed that these parallel fibers were difficult to distinguish from other components on R2 ratio maps, which hindered the reading of their degradation state. This highlights the importance of the cutting plane compared to the multifilament yarns direction.

Concerning the fibers that were perpendicular to the section plane, they apparently underwent bulk degradation as described in literature [START_REF] Hjort | Three-year results from a preclinical implantation study of a long-term resorbable surgical mesh with time-dependent mechanical characteristics[END_REF] and possibly followed the same degradation scheme as PLLA fibers since the inner part of the fibers seemed slightly more degraded. This observation however has to be taken carefully because fibers in TIGR meshes are much thinner than those in PLLA samples and resolution of the mapping may limit the accuracy of information provided on the crystallinity gradient. Moreover, amide I protein mapping revealed an important cell and ECM infiltration within the multifilament yarn structure between the fibers, higher than with PLLA samples after the same implantation time which could be explained by the thinner fibers and the textile structure of this device whose higher porosity stimulated cell proliferation and tissue colonization. Hence, the presence of ECM protein between the fibers may have induced misinterpretations from R2 ratio maps and no conclusion about auto-catalyzed degradation should be drawn.

Fig. 6

FTIR microscopy of non-pre-degraded TIGR mesh ( 0 TIGR) and slightly predegraded TIGR mesh ( 1 TIGR) explants, at 4 & 26 weeks, after subcutaneous implantation, in rabbits. See Fig. 5 for the legends http://eproofing.springer.com/journals/printpage.php?token=VY4BC44hemJOeae0hXG-KycQ870QGzyiLrSahy9V_DkqjTYsA8ScZw 20/27

The analysis of FTIR mappings of pre-degraded 1 TIGR 4 weeks after implantation led to the same observations as for 0 TIGR. Slow resorbing fibers were indeed still visible on OM pictures and PCA maps, apparently followed a bulk degradation process and seemed to be at the same degradation stage, probably because the implantation time was too short to have an actual stronger impact on the degradation state. ECM deposition within the textile structure was also observed, which should once again prompt to remain cautious about drawing conclusions from R2 ratio maps.

After 26 weeks of in vivo degradation, slow resorbing fibers of 1 TIGR samples were still visible but more difficult to distinguish from the surrounding tissues. Even though it was still possible to discern the shape of the fiber bundle on the PCA map, the fibers seemed smaller and more dispersed. On the opposite, the structure visible on the R2 ratio map appeared denser with larger stains than with PCA mapping. This difference could be explained by the important ECM infiltration around the fibers observed by protein mapping. Once again, the combined observation of the different mapping techniques was necessary to obtain reliable information. In this case, it was still possible to locate the fibers but not to have a precise idea of their degradation state because of the presence of neo-tissue ECM which blurred the polymer signal on the R2 ratio map.

Likewise, no reliable picture could be obtained from even more degraded samples (data not shown) because of high polymer fragmentation and tissue infiltration.

Discussion

The aim of this study was to evaluate the potential of a novel imaging technique for identification of the implanted material and the degradation monitoring of bioresorbable polymers with special attention paid to PLLA. This technique is mainly based on the progressive changes which were observed on FTIR spectra of PLLA (at 1209 and 1130 cm -1 ) during in vitro degradation and were attributed to the evolution of α crystal formation, resulting of the structure rearrangement from chain scissions and increased mobility. Firstly, R1 and R2 peaks intensity ratios (1209/1180 and 1130/1080 cm -1 respectively) were used to verify this hypothesis. Non degraded PLLA samples with identical chemical compositions but different crystallinity rates were obtained by applying different thermal treatments (both samples were melt, the first one was then slowly cooled to allow crystallization whereas the other one was quenched in liquid nitrogen to obtain an amorphous polymer sample). When used on these samples, R1 and R2 ratios calculation exhibited the same trend as observed in the in vitro degradation study. In both cases, an increase of these ratios was indeed observed with the crystallinity rate, which proved that (i) the change of R1 and R2 ratios values reflected the degradation state of the polymer samples and (ii) the evolution of these values resulted from changes in crystallinity rate and not in chemical composition. This was the first step to demonstrate the potential of this technique prior to working on actual explanted devices.

Mappings were then realized on PLLA samples which were submitted to different in vitro and in vivo degradation times. A combined analysis of OM pictures and FTIR mappings (PCA, R2 ratio, and amide I protein), obtained by the single Infra-Red microscopic observations of tissue sections allowed us to clearly identify the implanted material even at a fragmented state and to monitor the state of the PLLA fibers after implantation (only R2 ratio was selected to prevent interactions with protein peaks). It appeared that OM pictures used alone provided only limited information for this purpose, including location of the fibers and potential observation of their fragmentation in best cases. On the opposite, the association of the three FTIR mapping techniques led to further more precise interpretations. PCA mapping allowed much more precise location and visualization of the fibers in the tissue sections. R2 ratio mapping provided information on the crystallinity change of the PLLA fibers, which is directly related to the progress of their degradation state. Resolution of the maps was sufficient to observe degradation gradients within single fibers, showing that they had undergone faster inner degradation, which was consistent with the described degradation mechanism involving bulk hydrolysis with autocatalysis [START_REF] Anderson | Biodegradation and biocompatibility of PLA and PLGA microspheres[END_REF]. Amide I protein mapping revealed the presence of ECM [START_REF] Liu | Modification of the extracellular matrix following myocardial infarction monitored by FTIR spectroscopy[END_REF]. It is indicative of neo-tissue colonization of the implants, at least when it was detected within their structure, at advanced stages of degradation associated with implant fractures. Taken separately, each of these pictures thus provided information on the location of the PLLA fibers in the tissues, their fragmentation and degradation state and finally on the neo-tissue deposition. Nevertheless, complete and reliable interpretation required the analysis of these pictures taken all together. Pictures of the most degraded samples perfectly illustrated this necessity since they showed that an extensive presence of ECM proteins would interfere with the polymer peaks and make wrong signals appear on R2 ratio maps. Even though monitoring of the PLLA device was possible and reliable during a large period, a cautious study of the three different FTIR mappings together was essential to determine the limit time point up to which this technique should be used and trusted.

In order to see if this technique could be applied to other bioresorbable polymers, namely copolymers of lactic acid and other polyesters, a similar in vivo degradation protocol was used on another medical device, the TIGR matrix which is composed of two parts: slow resorbing fibers (copolymers of LA, GA and TMC) and fast resorbing fibers (copolymer of GA and TMC). In this specific case of multi-material device, characterization by FTIR mapping could be even more interesting since PCA mapping enabled the different components (ECM, fast resorbing and slow resorbing fibers) to be clearly identified at early degradation stages. It was thus possible to observe the morphology of the two kinds of fibers separately as long as the fast resorbing fibers were present. Moreover, the presence of LA in the slow resorbing fibers allowed the monitoring of their crystallinity and degradation state using R2 ratio mapping and evolution, just like it was done for PLLA samples. A precise evaluation of the degradation was however more difficult to obtain than with PLLA because of the low thickness of the fibers compared to the resolution of the mapping and because the textile structure allowed higher tissue deposition whose ECM proteins generated false signals on R2 ratio maps. Like for PLLA, FTIR mapping thus appeared to be a useful tool to follow the degradation and the integration of TIGR meshes after implantation, by following up the changes of the FTIR spectrum and specific peaks. Once again, only a careful FTIR combination of the different components, taken together (i.e. ECM, fast resorbing and slow resorbing fibers), but obtained from a single FTIR microscopy view, could lead to reliable information and interpretations.

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Conclusion

This study showed the value of an original technique to evaluate the evolution of resorbable polymers-based medical devices after implantation which is mandatory for any new commercial clinical implants. The novelty of this work lies in the association of the degradation state mapping of the polymers (based on the measure of crystallinity using specific FTIR peaks) with other already existing FTIR mappings (PCA and amide I protein peak) to get various valuable information: location of the materials, visualization of their morphology, comparative evaluation of their degradation state and assessment of their integration within the host tissues. This technique could be a very useful complement to current characterizations such as histological examination to judge the behavior of a resorbable medical device and the tissue response. If this concept was here demonstrated with PLA-based polymers, future developments on other resorbable polymers such as copolymers of lactic acid could be considered, as well as improvements of the process in order to reach semiquantitative measurements of the polymers degradation [37].

AQ2

Electronic supplementary material

The online version of this article (doi:10.1007/s10856-017-5894-7) contains supplementary material, which is available to authorized users.

Fig. 4 FTIR 2 27

 42 Fig.4FTIR microscopy of highly pre-degraded PLLA meshes ( 2 PLLA) explants, at 26 weeks, after subcutaneous implantation, in rabbits. Legends are the same as for Fig.2

Fig. 5 FTIR

 5 Fig.5FTIR microscopy of non-pre-degraded TIGR mesh ( 0 TIGR) explants, at 4 weeks, after subcutaneous implantation, in rabbits. PCA1 & PCA2 enabled the differentiation of slow resorbing yarns (copolymer of TMC, lactic acid; small diameter yarns, colored in orange or red) from fast resorbing yarns (copolymer of TMC, lactic acid and glycolic acid; larger diameter yarns, colored in blue-PCA1 -or green-PCA2). The intensities of R1 & R2 ratios and the amide I protein band is represented in false colors with a gradient from dark blue (low intensity) to hot red (high intensity). OM Optical microscopy, PCA1 & PCA2 Principal Component Analysis 1 & 2; R1 ratio: 1209/ 1180 cm -1 band intensity; R2 ratio: 1130/ 1080 cm -1 band intensity; Amide I protein band at 1646 cm-1

Table 1

 1 Properties of polymer samples used for in vivo degradation study before implantation

	Sample	Reference	Equivalent degradation time at 37 °C (weeks)	Mass retention (%)
	Non pre-degraded PLLA mesh	0 PLLA	0	100
	Partially pre-degraded PLLA mesh	1 PLLA	60	100
	Highly pre-degraded PLLA mesh	2 PLLA	120	59
	Non pre-degraded TIGR mesh	0 TIGR	0	100
	Partially pre-degraded TIGR mesh	1 TIGR	12	76
	Highly pre-degraded TIGR mesh	2 TIGR	80	43
	2.5. In vivo degradation study
	2.5.1. Animals
	In order to study the integration and the degradation of bioresorbable medical
	devices, an in vivo study was carried out. 18 New Zealand White male rabbits
	(Oryctolagus cuniculus; Charles River France) (3.0-3.8 kg) were included in this
	study.
	The study was performed in strict accordance with European Directive
	2010/63/UE and European Convention of the Council of Europe ETS123). All
	procedures were performed at NAMSA, (Chasse/Rhône, France).
	This study was approved by the local NAMSA Ethical Committee (NAMSA is
	an accredited facility registered with the French Department of Agriculture for
	animal housing, care and investigations).
	2.5.2. Experimental design

Table 2

 2

, Fig. 1 and Supplementary Fig 1).

Table 2

 2

	Degradation time (hours)	0	8	24	120
	Mn retention (%)	100	60	15	< 5
	Mass retention (%)	100	> 99	> 97.5	70
	Crystallinity (%)	48	54	57	73
	Eq. degradation time at 37 °C in weeks	0	23	71	354

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Characteristics of in vitro degraded PLLA meshes

Table 3

 3 R1 and R2 ratios values of PLLA samples, calculated from FTIR spectra. Ratios R1 (1209/ 1180 cm -1 band intensity) and R2 (1130/1080 cm -1 band intensity) were selected from FTIR spectra as indicators of PLLA cristallinity

	Polymer	R1	R2
	Amorphous PLLA	0.06	0.19
	Crystallized PLLA	0.12	0.28
	Non-degraded PLLA	0.11	0.28
	highly degraded PLLA	0.41	0.49
	3.2. In vivo degradation of PLLA meshes

(See Table

Acknowledgements

This work was partially supported by the Rhône-Alpes Region and the Fonds Unique Interministériel program (No AAP:15-Project No F1312025 V REVAMED).

Compliance with ethical standards

Conflicts of interest

The authors declare that they have no competing interests.

Electronic supplementary material

Supplementary Data

Supplementary Figure 1 Supplementary Figure 2


Anthony Lannes

Yan Suffren

Jean Bernard Tommasino

Rodica Chiriac

François Toche

Lhoussain Khrouz

Florian Molton

Carole Duboc

Isabelle Kieffer

Jean-Louis Hazemann

Lhoussain Khrouz∫

Florian Molton¦

Carole Duboc¦

Isabelle Kieffer¤

Jean-Louis Hazemann¤

Christian Reber

Andreas Hauser

Dominique Luneau

Room Temperature Magnetic Switchability Assisted by Hysteretic Valence Tautomerism in a Layered Two-Dimensional Manganese-Radical Coordination Framework

Keywords: manganese, nitronyl-nitroxide, valence tautomerism, magnetism, solid-state absorption, electronic paramagnetic resonance, X-ray spectroscopy

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

INTRODUCTION

Developing bistable systems responding to external stimuli such as heat, magnetic fields or light is of high scientific interest and important for technological purposes. Such systems are relevant for novel applications in switches, sensors, signal processing and information storage.1-8 Potential candidates are valence tautomeric (VT) complexes, consisting of a redox-active metal center coordinated to non-innocent ligands.9-11 VT transitions have their origin in the capacity of switching the metal-to-ligand electronic charge distribution reversibly between two redox isomers. The first example was reported by Pierpont et al. in 1980, in a cobalt(III) complex bearing a catecholate (CAT) ligand. 12 In this redox process, the Co(III) ion is reduced to Co(II) and the CAT ligand is oxidized to its semiquinonato (SQ) form. VT transitions involving SQ-CAT ligand are by far the most reported in the literature, mainly in Co,12-27 and Mn isolated complexes,28-33 but also in various d block elements such as Fe, Ni, Rh, and V.34-39 VT transitions were also observed in four unidimensional (1D) and one Co-based bidimensional (2D) coordination polymers. 32,[40][41][42][43] Other ligands than SQ-CAT are less common in VT systems, with a dozen examples reported so far, involving porphyrin derivatives, cyclopentadienyl, or bis(diisopropyl-phenylimino)acenaphthene.44-55In the latter case VT was reported for Yb. 55 Nitroxide radicals are potentially redox-active and may be reversibly oxidized to an oxoammonium cation, or reduced to an aminoxyl anion (Scheme 1).56 However, they have generally been considered to be stable, even when coordinated to metal ions. This is especially true for the nitronyl-nitroxide (NIT) radicals where the radical is delocalized across both NO groups.57 Scheme 1. Redox isomers of the nitroxide radical.

The stability of NIT radicals, combined with their bridging ability, has greatly contributed to their extensive use as spin carriers for the efficient building of molecule based magnets, based on the so-called metal radical approach.58 Exploiting this approach, we have developed chelating NITs especially with imidazole (NITImH) and benzimidazole (NITBzImH) substituents,59 for the synthesis of discrete complexes with d60 and f61-65metal ions as well as 1D66 and 2D coordi-nation67,68 polymers. The 2D polymer {[Mn2(NITBzIm)3]ClO4}n, obtained with NITBzImH and manganese(II) perchlorate can be described as a layered compound and behaves as a ferrimagnet with a Curie temperature Tc = 55 K, the highest reported so far for NIT-based compounds. With the aim of fully understanding the origin of such a dramatic change in magnetism, we report here an in-depth investigation of {[Mn2(NITIm)3]ClO4}n (1), combining a barrage of complementary techniques based on solid state temperature-dependent studies including crystal structure determination, magnetic susceptibility measurements, differential scanning calorimetry (DSC), electron paramagnetic resonance (EPR), UV-Vis spectroscopy, X-ray absorption (XAS) and emission (XES) spectroscopies of this 2D coordination polymer. The overall data can be rationalized by a solid state valence tautomerism in the room temperature domain. The transition is highly cooperative with a thermal hyste-resis, and is only observed in the solid state. Compound 1 represents the first example of a valence tautomerism transition involving nitronyl-nitroxide radicals, and the first example occurring in a Mn-based 2D coordination polymer.

RESULTS AND DISCUSSION

Magnetic behavior. Figure 1 shows the temperature dependent magnetic susceptibility measurements of 1, in the form of χT vs. T under a magnetic field of 1000 Oe and in the 375value is 7.20 emu•K•mol--1 at 280 K then drops abruptly with an inflexion point at 274 K to reach a minimum of 5.81 emu•K•mol-1 at 250 K. On further cooling, value of 8.23 emu•K•mol-1 at 5 K. On heating the sample back curve is well superimposed on the cooling curve from 5 K to 250 K. The heating curve keeps decreasing until 280 K where it abruptly increases to reach a value of 7.93 emu•K•mol-1 at 302 K with an inflexion point at 294 K. The heating curve decreases to 7.26 emu•K•mol-1 at 375 K, again superimposed to the cooling curve. This magnetic behavior is generally indicative of a first order phase transition with a hysteresis of 20 K. The 1/ vs. T curves on either side of the transition follow a Curie-Weiss law (Figure S1) with identical slopes corresponding to a Curie constant (C) of 5.27 cm3•K•mol-1, suggesting that the total spin of 1 at high temperature (1HT) is This is in agreement with ferrimagnetic behaviors on both sides of the step, as previously reported for 1D and 2D Mn(II)-NITIm and Mn(II)-NITBzIm compounds.60,68 This ferrimagnetic behavior is expected, at least for 1HT, due to the strong antiferromagnetic coupling between the Mn(II) ions and the bridging radicals.66 It is remarkable that the Curie ) constants of 1/ vs. T in the HT regime are close to those found for the 2D system with NITBzIm, that is, {[Mn2(NITBzIm)3]ClO4}n (C = 5.33 cm3•K•mol-. The same holds for (Figure S2). 68 Assuming that 1HT corresponds to [MnII2(NITIm)3ClO4]n, the total ground-state spin may be approximated to S = 7/2 -per unit formula. Because the Weiss constant dramatically decreases below the transition temperature, but the slopes do not change, this behavior was first ascribed to a weakening of the magnetic coupling between the Mn(II) ions and the bridging radicals at low temperature. This weakening could be due to some flexibility in the coordination sphere that results in a decrease in the overlap of the Mn(II) and radical magnetic orbitals. This explanation was increasing upon cooling similar as for the Mn(II)-NITBzIm system. Moreover, ZFC-FC studies at low temperature evidence a remanent magnetization in agreement with a ferrimagnetic behavior below 1.4 K. 67 In Figure 1, -1 at 5 K and gradual decrease to 5.81 emu•K•mol-1 is due to saturation effects under the magnetic field of 1000 Oe. Under a lower magincreasing at low temperature (Figure S3). The magnetization at 2 K approaches saturation netically coupled to 3 radicals (Figure S4). This low value could originate from canting, as also observed for the 2D system with NITBzIm {[Mn2(NITBzIm)3]ClO4}n.68 However, investigating compound 1 in more detail with a set of different techniques, we have come to a better explanation, namely that the transition corresponds to valence tautomerism: two Mn(II) of 1HT are both oxidized to Mn(III) in 1LT, while two among the three NITIm-radicals are reduced to their diamagnetic analogs so-called NITRed here (Scheme 2) to give the compound {[Mn2(NITIm)(NITRed)2]ClO4}n. The resulting system can also be approximated to a total ground-state spin of 7/2 in agreement with the identical slopes of the side of the transition. The increase is thus explained by the delocalization of the radical.

Differential Scanning Calorimetry.

The DSC shows an exothermic peak at 274 K upon cooling and an endothermic peak at 287 K upon heating (Figure 2). The peak temperature on cooling is in agreement with the transition temperature found by the magnetic measurements (see above) and by absorption and EPR spectroscopies (see below), while the peak temperature on -1 -1•mol--12.08 kJ•mol-1 and S = -42.08 J•K-1•mol-1 upon heating. The heat capacity measured by DSC with the Cp of sapphire as reference (Figure S5) decreases smoothly upon cooling with a broad peak at 274K and increases continuously on heating with a broad peak at 287K. We further performed modulated-temperature DSC measurements (Figures S6-S7).69 Whether upon cooling or heating, they both exhibit the reversing and the non-reversing components which respectively discriminate the sensible and the latent heats, and unambiguously demonstrates that the transition is of the first order type.70,71 More surprisingly, we found that the nonreversing peak (latent heat) splits at low heating rate. Thus at 2 K•min-1 (Figure S7) there are a main broad peak at 285 K and a second broad flattening peak centred at 300 K, which roughly correspond to the beginning and the end of transition, as seen on heating from the magnetic measurement (Figure 1). This effect is not seen upon cooling. Such a splitting and dependence on heating rate is reminiscent of some previous studies on solids exhibiting transition upon disproportionation.71-73 We believe this may be paralleled with the valence tautomerism at work in our compound. Following we ascribe the two latent heat peaks (285K and 300K) to the non-simultaneity of the electron transfer involving the two crystallographically independent manganese ions.

X-ray single crystal structure analysis. The crystal structure of 1 has been reported previously at 313 K (1HT) and 193 K (1LT).67 X-ray data have also been recorded at 300 K and 233 K, in order to get additional insight into the temperaturedependency of the structural properties of 1. At all temperatures the structure could be refined in the same space group, namely P21, indicating that no structural phase transition occurs. The crystal structure parameters are listed in Tables 1 andS1-S3. The variation of all parameters is less than 4 %. The evolution of metal-ligand bond lengths varies from bond to bond (Table S2). The room temperature X-ray powder pattern is similar to the calculated ones at 313 K and 300 K, with significant changes compared to calculated ones at 233 K and 193 K (Figure S8).

The crystal structure determination of 1 at 300 K confirms the previously reported structure at 313 K.67 The asymmetric unit is composed of two independent Mn centers, coordinated to three bidentate NITIm-anion radicals in a distorted octahedral geometry (Figure S9 ycomb-like structure, organized as sheets of formula [Mn2(NITIm)3]+. The sheets are arranged with the Mn(II) centers on top of each other along the a direction and with the ClO4-anions in the voids between the sheets to form a layered structure (Figures 3 andS5-S6).

When decreasing the temperature, the overall layered structure is kept in line with the unchanged space group. However, examination of the coordination sphere reveals dramatic changes as shown in Table 2. At 313 K, the Mn-O and Mn-N bond lengths are homogenous between 2.16-2.23 Å with a small shortening at 300 K, which can be ascribed to temperature effects. These bond lengths are typical for Mn(II) coordination with N-and O-based ligands, particularly with Mn-nitroxide compounds.60,66

At 233 K and 193 K the picture is notably different. The range of Mn-O and Mn-N bond lengths is no longer homogenous, with values ranging from 1.83-2.24 Å at 233 K and 1.85-2.27 Å at 193 K, with four long and two short bonds. On the one hand, the short distances (1.83 Å and 1.85 Å) are unrealistic for Mn(II). On the other hand, the overall contraction along one preferential axis is the signature of a compressed Jahn-Teller Mn(III) system.74-76 This is substantiated by the Bond Valence Sums calculated for each manganese ion at each temperature (Table S2).77They reveal that both manganese ions are closer to the +II oxidation state at temperatures above the transition (313 K and 300 K), and closer to the +III oxidation state at lower temperatures (233 K and 193 K).

The N-O bond lengths of the NITIm-moiety are summarized at each temperature in Table S3. On the one hand, those usually encountered in NIT radicals coordinated to a manganese ion are in the 1. 27 C) are affected by the elongation of bond lengths while the third one (namely B) displays classical N-O• bonds over the temperature range in agreement with our above description of VT process in 1. However, a close analysis shows that it is difficult to unambiguously identify which radicals are reduced or not, as short and long bonds are found on a same "NITIm" moiety and the N-O mean bond length is 1.33(3) Å, even at 193 K. This suggests that in the low temperature region the remaining radical is delocalized on the three "NITIm" moieties. Such a delocalized mixed valence state for the radical is in agreement with the long range magnetic order observed at low temperature. Moreover, the Jahn-Teller distortion of the Mn(III) metal ions may be the cause of the 20K width hysteresis loop because it is known to stress cooperative effects in different class of compounds such as the closely related spin crossover systems.84-86 UV-vis spectroscopy. To confirm the aforementioned hypothesis, temperature-dependent solid-state UV-Vis spectra were carried out. Figure 4 shows the absorption spectra in heating mode (Figure S10 for cooling mode). At 310 K the spectrum shows a broad and intense band with two maxima in the range 35000-23500 cm-1 (290-425 nm), corresponding to π  π* and n  π* transitions from the NITIm-ligands, respectively. Another broad band centered at 12660 cm-1 (790 nm) is visible at 310 K.

On cooling from 310 to 250 K, the maximum of this band slightly shifts to the blue, concomitant with a decrease of the integrated intensity by around 50 %. The bands in the 35000 to 23500 cm-1 range likewise decrease in intensity. Concomitantly, a new band appears in the range 23500 to 15400 cm-1 (425 -650 nm) and its intensity increases as temperature decreases. Upon subsequent heating from 250 K back to 310 K the evolution of the absorption spectrum is fully reversible. The temperature-dependent quantitative evolution of the absorbance at 510 nm, where the most pronounced difference occurs, is depicted in Figure 5 in the interval from 250 K to 310 K. It shows that the optical and magnetic data, on cooling and on heating, are perfectly superimposable, reproducing the 20 K hysteresis. Thus the changes in the absorption spectrum as a function of temperature are clearly associated with the phase transition and show that the associated changes in the electronic structure are quite dramatic. The Raman spectra excited at 633 nm and 785 nm show clear peaks in this region. The most intense peak of this group shifts from 1417 cm-1 at room temperature to 1421 cm-1 at low temperature (Figure S11). Such small changes cannot be unambiguously attributed to VT transition as all frequencies are expected to increase at low temperature.90 EPR spectroscopy: Powder X-band EPR spectra of [Mn2(NITIm)3ClO4]n 1 have been recorded between 310 and 110 K, and displays a unique isotropic feature centered around g = 2 (Figure 7a). Unexpectedly, the signal intensity decreases as temperature decreases. A drastic decrease of the intensity is observed on cooling around 276 K, while the drastic increase occurs around 296 K on heating. This temperature behavior parallels that of the magnetic susceptibility and solid-state absorption spectroscopy, evidencing that the isotropic EPR signal is the signature of 1 at high temperature.

The 20 K hysteresis can be reproduces with a double integration of the EPR signal (Figure 7b). Interestingly, at low temperature and in parallel mode X-band EPR, a low-field EPR signal is observed, characteristic of high spin Mn(III) (Figure 8).91,92 This is consistent with the fact that below the transition temperature, the oxidation state of the Mn ions can be assigned to +III. EPR is thus in overall agreement with the other data, and provides a direct evidence of the Mn oxidation state in 1LT. Mn K edge X-ray emission spectroscopy (XES). XES and XAS measurements have been carried out on the CRG-FAME beamline (BM30B) at the ESRF (Grenoble, France).93 XES spectra were recorded at 300 K and 10 K for 1 and for two Mn(II) and Mn(III) reference compounds (Figure 9). The Mn At 300K the emission data show that the spectrum of 1HT is superimposed with the spectrum of the Mn(II) reference68 while at 10 K the spectrum of 1LT is very similar to the one of the Mn(III) reference. Intensity ratios are identical. Kβ' line positions are at 6.4775 keV for the references and for 1HT and 1LT.

Kβ1,3 maximum is at 6.4230 keV for Mn(II) reference and 1HT at 300 K and is shifted towards the low energies by 0.5 eV at 6.4225 keV for Mn(III) reference and 1LT at 10 K. This is a good indication for the presence of Mn(II) at HT and Mn(III) at LT, consistent with the VT transition.

Mn K edge X-ray absorption spectroscopy (XAS). The XANES spectra (Figure 10a) can give information on the Mn valence through two characteristic features: the pre-edge signal and the edge position. At 300 K, the pre-edge signal of 1HT has the same shape and position as the pre-edge signal of the Mn(II) reference (6.5405 keV), irrespective of whether the spectrum of 1HT is collected at the Mn(II) or Mn(III) maximum energy emission, showing that the precise energy of collection has no influence on the final result (Figure 10b). The edge positions of 1HT and of the Mn(II) reference is at the same energy (6.546 keV), confirming the +II oxidation state for 1HT (Figure 10c). At 10 K, the pre-edge signal of 1LT has the same shape and position as the pre-edge signal of the Mn(III) reference and is at the same energy as the HT spectrum (6.5405 keV), but the intensity is lower. Compared to the Mn(II) reference, the edge positions of the Mn(III) reference and 1LT collected at the Mn(III) maximum energy emission are shifted toward higher energy by ca. 2.5 eV at 6.5485 keV. This indicates a transition from Mn(II) to Mn(III) at low temperature.

1LT evolves under the beam, and the edge position is shifted towards lower energies (0.1 eV•h-1). Thus, only the first spectrum is presented here, explaining the poorer statistics compared to the five other spectra. The difference of 1 eV between the 1LT spectra, collected at the Mn(III) emission maximum and at the Mn(II) emission maximum, is mainly due to this evolution under the beam, since the latter was collected at the end of the experiment. Nevertheless, the spectra confirm the consistency of the results whatever the energy of collection.

CONCLUSION

In this paper we report the detailed solid-state study of the thermo-induced hysteretic transition exhibited by the 2D coordination polymer {[Mn2(NITIm)3]ClO4}n. Temperature-dependent studies using a variety of techniques including single crystal X-ray crystallography, magnetic susceptibility measurements, DSC, EPR, UV-vis absorption spectroscopy, XAS and XES have been performed and all demonstrate that this thermo-induced transition is unambiguously due to a valence tautomerism process. Along this VT process the two manganese ions in their +II oxidation state at high temperature are oxidized to the +III oxidation state at low temperature and two of the three NITImH radicals are reduced to the NITRed aminoxyl and diamagnetic anion (Scheme 2). From electrochemistry experiments (Supporting Information) and in agreement with previous reports94 the E1/2 redox potential aminoxyl anion/nitroxide radical, is -0.25V/SHE (Standard Hydrogen Electrode) as shown in Figure S12 while the redox potential of the system Mn3+/Mn2+ is given higher at 1.54 V/SHE. This is in favor for stable MnII-nitroxide complexes as it is generally observed. Thus, the simple comparison of the redox potential cannot explain this thermo-induced VT. This may be due that the redox potential are determinate in solution while VT happens here in the solid state. Multiscale quantum calculations that are in progress should give a better view.

To the best of our knowledge this is the first metal-nitroxide radical system to exhibit bistability based on valencetautomerism. Moreover, this is associated with a 20 K hysteresis at room temperature. This opens novel perspectives for the so called metal-radical approach. Indeed, beyond the interest of nitronyl-nitroxide radicals as spin-carriers bridging ligands, this result shows that their redox activity can give rise in some cases to metal-radical electron transfer. This may be an additional advantage in the context of multifunctional materials, especially in the search for switchable magnets. The compound reported here is based on layered 2D metal-radical coordination polymer. It is noteworthy that the [Mn2(NITIm)3]+ layers are cationic with the intercalated perchlorate anions in between, to afford a layered 3D chiral structure. From this point of view, it is a multifunctional material that gathers many interesting functionalities such as room temperature bistability, ferrimagnetic behavior and chirality. The synthesis is general and may be extended to the related imidazole type substituted nitronyl-nitroxide radical while almost any type of anions can be intercalated. This allows a great versatility to tune the temperature of the transition or implement the multifunctionality, and accordingly work is in progress in the group.

ASSOCIATED CONTENT

Experimental details, complete crystallographic data and figures giving additional data. This material is available free of charge via the internet at http://pubs.acs.org.

67 Scheme 2 .

 672 Scheme 2. Redox isomers for NITImH free radical.

Figure 1 .

 1 Figure 1. Temperature dependence of χmT for 1, in a 1000 Oe magnetic field upon cooling (•) and heating (•).

Figure 2 .

 2 Figure 2. Differential scanning calorimetry (DSC) for 1 upon cooling (▼) and heating (▲) at 10 K•min-1.

Figure 3 .

 3 Figure 3. a) Crystal packing in the bc plane showing bidimensional honeycomb-like structure at 300 K. b) Crystal packing in the ab plane showing layers piling up along a axis and interlayer space filled with ClO4-anion at 300 K. Hydrogen atoms have been removed for clarity.

Figure 4 .

 4 Figure 4. Temperature dependence of the solid-state absorption spectra for 1 during heating. Temperature rate was monitored at 0.25 K•min-1.

Figure 5 .

 5 Figure 5. Superimposition of 510 nm absorption (circles) with magnetic susceptibility (squares) during cooling (blue) and heating (red). Lines are eye-guides.

Figure 6 .

 6 Figure 6. Variable-temperature Raman spectra of 1 at three excitation wavelengths.

Figure 7 .

 7 Figure 7. (a) Temperature dependence of the solid-state X-band EPR spectra for 1 during heating. (b) Superimposition of the magnetic data (square) and the intensity of the EPR lines (triangle) during cooling (blue shades) and heating (red shades). Lines are eye-guides.

Figure 8 .

 8 Figure 8. Powder X-band EPR spectrum of 1 recorded at 21 K in parallel mode.

  (II) reference corresponds to the 2D {[Mn2(NITBzIm)3]ClO4}n compound and the Mn(III) reference corresponds to a [Mn(ImIm)2(CH3OH)2]ClO4 isolated complex (ImIm-= 2-(2-imidazolyl)-4,4,5,5-tetramethyl-4,5-dihydro-1H-1-oxyl).

Figure 9 .

 9 Figure 9. X-ray Emission Spectra (XES) for 1 collected at high and low temperatures for comparison with the Mn(II) and Mn(III) references.

Figure 10 .

 10 Figure 10. a) XANES Spectra (XES) for 1 collected at high and low temperatures with the Mn(II) and Mn(III) references. b) Zoom on the pre-edge signal area. c) Zoom on the edge position area. Brown and black dotted lines are maximum intensity for the Mn(II) and Mn(III) references, respectively.

Table 1 .

 1 Crystal data of 1 at several different temperatures.

  -1.33 Å range, regardless of the oxidation state.78-82 On the other hand, the N-O-single bond is close to 1.40 Å, while the N=O double bond is close to 1.21 Å.83 For 1, at 313 K, the N-O bond lengths are quite homogenous with a mean value of 1.29 Å (2) well in agreement with a nitroxide radical. At lower temperatures and particularly at 233 K and 193 K, the N-O bonds length show a higher spreading with some close to 1.40 Å consistent with N-O-single bonds. A quick view of the N-O distances may suggest that only two radicals (namely A and

Table 2 .

 2 Mn-Ligand Bond lengths for 1 at several temperatures.

	Bonds (Å)		313 K	300 K	233 K	193 K
	Mn1	N3A	2.22(2)	2.19(3)	2.19(3)	2.15(3)
	Mn1	N3B	2.26(2)	2.24(2)	2.10(3)	2.07(2)
	Mn1	N3C	2.11(3)	2.14(2)	2.11(3)	2.07(3)
	Mn1	O1A	2.18(2)	2.18(2)	1.93(3)	1.90(3)
	Mn1	O1B	2.29(2)	2.28(2)	2.26(3)	2.27(2)
	Mn1	O1C	2.16(2)	2.07(2)	1.83(3)	1.95(2)
	Mn2	N4A	2.17(2)	2.08(2)	2.06(4)	2.08(3)
	Mn2	N4B	2.17(3)	2.14(2)	1.95(4)	2.07(3)
	Mn2	N4C	2.20(2)	2.19(2)	1.94(3)	1.99(3)
	Mn2	O2A	2.19(2)	2.15(2)	1.89(3)	1.85(2)
	Mn2	O2B	2.20(2)	2.16(2)	2.24(2)	2.22(2)
	Mn2	O2C	2.23(2)	2.18(2)	2.05(2)	2.04(2)

ACKNOWLEDGMENTS

A.L. held a CMIRA Explora' doc 2012 fellowship from la Région Rhône-Alpes. D.L. and C.R. are grateful for funding by the Centre National de la Recherche Scientifique (CNRS) for collaborative research (PICS 2015-2017). The authors are also grateful to the EPR facilities available at the national EPR network (IT CNRS 3443).

AUTHOR INFORMATION

Corresponding Author * Corresponding authors: dominique.luneau@univ-lyon1.fr.

Notes

Authors declare no competing financial interests.