Geoffrey Ginoux
email: geoffrey.ginoux@univ-reims.fr

Philippe Dony

Isabelle Vroman

Sebastien Alix

IMPROVING THERMOMECHANICAL PROPERTIES OF FUSED FILAMENT FABRICATION PRINTED PARTS BY USING NANOCOMPOSITES

Keywords: A. Polymer-matrix composites (PMCs), B. Microstructures, B. Thermomechanical, C. Statistical properties/methods, Additive manufacturing

This study aims to assess the viability of nanocomposites in fused filament fabrication. Polylactide/organo-modified layered silicate materials have been studied in terms of structure and properties evolution through the whole manufacturing chain, and for thermomechanical optimisation. The method included the optimisation by the composite formulation and by the additive manufacturing process by experimental designs. A competition between flow thinning of the molten composite and stiffness of the input filament with the addition of filler allows an enhanced printability at intermediate loadings. The method provides an optimisation of the final parts, enabling similar or greater tensile properties, while providing better thermal stability, with additively manufactured nanocomposite materials compared to conventionally manufactured polymers.

Introduction

Initially used for rapid prototyping, additive manufacturing (AM) processes offer the possibility to manufacture complex parts thanks to their layer-by-layer assembly, which enables the fabrication of expensive end-use products at lower cost, or of impossible-to-manufacture parts with usual subtractive or formative techniques. However, a loss in mechanical performance is observed with additively manufactured parts compared to those produced by subtractive and formative processes, which prevents the development of this technology towards small-series production or medical applications [START_REF] Ngo | Additive manufacturing (3D printing): A review of materials, methods, applications and challenges[END_REF]. This development is a major issue and requires new materials to design fully functional products [START_REF] Bekas | 3D printing to enable multifunctionality in polymer-based composites: A review[END_REF]. The addition of fillers in a polymer matrix is one way to adapt the rheological and functional properties, and to expand their potential uses [START_REF] Wang | 3D printing of polymer matrix composites: A review and prospective[END_REF]. Fused Filament Fabrication (FFF) technique is one of the most commonly AM process used due to its speed and costeffectiveness [START_REF] De Leon | High performance polymer nanocomposites for additive manufacturing applications[END_REF], but the use of micrometric fillers could cause the nozzle to clog because of the filler dimensions compared to the low nozzle diameter [START_REF] Campbell | 3D printing of multifunctional nanocomposites[END_REF]. Nanofillers appear as an obvious solution to get around this geometric problem. One particular nanofiller is the organo-modified layered silicate (OMLS), a clay (phyllosilicate) in the form of nanoplatelets of around 1 nm in thickness and 100 to 1000 nm in width and length. These dimensions develop a high specific surface area of around 800 m²/g, which allows the properties development at low loadings, less than 10 wt% against 40 to 60 wt% for microfillers [START_REF] Hassan | Polylactic Acid Based Blends, Composites and Nanocomposites[END_REF], which would again decrease the nozzle clogging risk. Ray et al. [START_REF] Ray | New polylactide/layered silicate nanocomposites. 5. Designing of materials with desired properties[END_REF] revealed several property improvements thanks to OMLS, from thermal stability and gas barrier properties, to mechanical reinforcement. Consequently, the use of nanocomposites is one promising solution to move the FFF technology towards small-series productions of functional end-use products [START_REF] Ivanova | Additive manufacturing (AM) and nanotechnology: promises and challenges[END_REF]. With the use of polylactide (PLA) as the matrix, a bio-sourced, biodegradable and biocompatible aliphatic polyester, the applications are expected, but not limited, in aerospace, medical field or packaging [START_REF] Armentano | Multifunctional nanostructured PLA materials for packaging and tissue engineering[END_REF]. Independently of the kind of fillers, a thickening effect is observed with the addition of particles in a molten polymer [START_REF] Mueller | The rheology of suspensions of solid particles[END_REF], which could prevent the materials to flow through the nozzle. The printability of nanocomposite systems is then not certain. An optimisation of the OMLS loadings is expected with a competition between the ability to be shaped by FFF and the development of thermomechanical properties. The aim of this article is to assess the viability of additively manufactured PLA/OMLS-based materials while simultaneously reaching similar or greater thermomechanical performance than formative neat PLA one. Indeed, it is wellknown that printed parts generally have lower mechanical properties compared to injected parts, and therefore have lower thermodynamic properties. Thermomechanical performance in this study includes thermal stability and tensile properties, namely Young's modulus, ultimate tensile strength and elongation at break. PLA/OMLS-based nanocomposites silicates are proposed in fused filament fabrication to extent this technology towards the small series production of functional and customisable parts. Targeted applications may include on-demand part production during long-duration space missions where high thermomechanical performance are needed, customisable packaging where barrier properties are sought, or medical applications for drug control delivery in bone tissue engineering. As the rheology is dependent on temperature, shear rate and OMLS loading [START_REF] Ginoux | Decrease in non-linear viscosity of a polylactide nanocomposite with regard to the clay volume fraction[END_REF], the optimisation must consider the thermokinetic process parameters in addition to the formulation. The viability of nanocomposites includes (1) the possibility of use and (2) the capacity of the technology to freely tailor thermomechanical properties. A method is introduced in this study, including multi-physics characterisation and modelling/statistical tools for performance tailoring. Thermal properties are characterised by thermogravimetric analysis (TGA) to obtain the thermal stability and by differential scanning calorimetry (DSC) to assess the crystallinity rate and the phase transitions. Mechanical properties are characterised by tensile tests and optimised using Doehlert designs, a response surface methodology for properties optimisation and process understanding with a time and cost efficiency [START_REF] Ginoux | Influence of fused filament fabrication parameters on tensile properties of polylactide/layered silicate nanocomposite using response surface methodology[END_REF]. Two factors are considered, namely printing speed and temperature. Flow behaviour is defined by capillary rheology. Formulation state is determined by gel permeation chromatography (GPC) for polymer molar mass, TGA for filler loading, and X-ray diffraction (XRD) and transmission electronic microscopy (TEM) for filler dispersion morphology.

Materials and methods

The overall experimental section is depicted in the following schematic diagram (see Figure 1).

2.1.

Materials and nanocomposite preparation PLA was purchased from NaturePlast (Ifs, France) in pellets form under the reference PLE 005, an extrusion grade. PLA's specific gravity is 1.26. The OMLS was an organically modified montmorillonite (MMT) supplied by Southern Clay Products (Gonzales, TX, USA) under the trade name Cloisite ® 30B (C30B). The modifier was the methyl, tallow, bis-2hydroxyethyl, and quaternary ammonium with a concentration of 90 mEq per 100 g. C30B's specific gravity is 1.98, and that of native MMT is 2.86. PLA was first dried at 70 °C overnight before nanocomposite compounding with a co-rotating twin-screw extruder Clextral BC21 (D = 25 mm, L/D = 36) with temperatures set between 165 °C and 180 °C at 300 rpm. The residence time was approximately three minutes with a mass flow of 1.8 kg/h. A masterbatch at 13.2 wt% of clay in PLA was obtained. The masterbatch was then dried and diluted with PLA at desired loadings of 0.5, 1, 3, 5, and 7 wt%. Samples were named PLA05, PLA1, PLA3, PLA5, and PLA7, respectively. PLA also underwent the extrusion process, and the sample was named PLA0. This method permitted us to reach a better nanoparticle dispersion and orientation, as described by Alix et al. [START_REF] Alix | Effect of highly exfoliated and oriented organoclays on the barrier properties of polyamide 6 based nanocomposites[END_REF] for polyamide nanocomposites, which is essential to develop the composite properties compared to the added amount of clay [START_REF] Paspali | The influence of nanostructure on the mechanical properties of 3D printed polylactide/nanoclay composites[END_REF]. Materials were then pelletised using a Scamia granulator 960001 GR. Pellets were first dried overnight at 70 °C in a ventilated oven and then shaped into a 1.65 ± 0.05 mm diameter filament thanks to a single screw extruder Axon BX10 (D = 10 mm, L/D = 18) with a screw speed at 20% of maximum speed and a 0,46 kg/h mass flow. The profile temperature was set to 190 °C, 185 °C, 185 °C from feeding zone to nozzle. Filaments were cooled with a water bath at room temperature. The filament was pulled by a Haake PP1 to obtain the required diameter.

Once filament obtained, tensile test specimens were manufactured using a Makerbot Replicator 2 with a 0.4 mm diameter nozzle. The specimens were modelled using Solidworks software with dimensions according to 5A-type specimens from ISO 527-2 standard. STL file was then imported to the Makerbot Print software. Specimens were printed with two shells, in a 90°/0° crossed-deposition fashion and with 0.2 mm layer height since these parameter values gave the optimal mechanical properties [START_REF] Li | Composite Modeling and Analysis for Fabrication of FDM Prototypes with Locally Controlled Properties[END_REF][START_REF] Sood | Parametric appraisal of mechanical property of fused deposition modelling processed parts[END_REF][START_REF] Yao | Tensile failure strength and separation angle of FDM 3D printing PLA material: Experimental and theoretical analyses[END_REF]. Infill percentage was set to a theoretical 100%. Printing speed and extrusion temperature were kept constant for the whole part and were defined according to the Doehlert method. A Doehlert network with two experimental designs, including ten experiments, was used to evaluate the tensile properties regarding printing speed from 15 to 105 mm/s in seven levels and extrusion temperature from 190 to 240 °C in three levels according to Ginoux et al. work [START_REF] Ginoux | Influence of fused filament fabrication parameters on tensile properties of polylactide/layered silicate nanocomposite using response surface methodology[END_REF]. One should note that PLA7 was not printed because of its rigidity. The PLA7 filament did not accept the strain applied by the roller driving the filament towards the printing head. The broken filament was then impossible to print with FFF and only PLA0 and nanocomposites loaded under 5 wt% were printed. To compare tensile properties of printed parts to fully infill parts manufactured by a formative process, materials in pellet form after compounding were injected moulded into 5A-type specimens with a Babyplast 6/10P equipped with an 18 mm diameter piston and a Moullo X Dry Air to preheat pellets at 70 °C and keep them moisture-free before injection moulding at 200 °C.

2.2.

Characterisation tests Before each measurement, all samples were dried at 80 °C under vacuum (-72 cmHg) for 3 h. TGA were performed for each formulation and after each process to assess the thermal stability regarding formulation and the remaining MMT mass fractions through successive extrusions. A TG209 F3 from Netzsch was used. Samples were heated from 25 °C to 800 °C at a 20 K/min heating rate under nitrogen flow of 3 mL/min. The MMT mass fraction was controlled by thermogravimetric analysis from residue after degradation at 800 °C for 10 min under air flow for pellets, filaments, and tensile test specimens. Each measurement is repeated three times. Calorimetric measurements were performed on filaments to assess the behaviour of nanocomposites during melting and their temperature of use. For this purpose, only the first heating is presented in this study. A DSC204 F1 from Netzsch was used and calibrated with Indium standard. Samples were analysed from 25 to 210 °C at a 10 K/min heating rate three times for standard deviation calculation. Equation (1) was used to calculate the crystallinity rate, 𝜒 𝑐 .

𝜒 𝑐 = Δ𝐻 𝑚 -Δ𝐻 𝑐𝑐 Δ𝐻 𝑚 0 × 1 1 -𝛷 𝑚 ( 1
)
Δ𝐻 𝑚 is the measured melting enthalpy, Δ𝐻 𝑐𝑐 the measured cold crystallisation enthalpy, Δ𝐻 𝑚 0 the melting enthalpy of a theoretical 100% crystalline PLA (93 J/g [START_REF] Dong | Carbonized cellulose nanofibers as dielectric heat sources for microwave annealing 3D printed PLA composite[END_REF]), and 𝛷 𝑚 the MMT mass fraction. PLA molar masses were determined by GPC for each formulation after compounding to determine the molar mass degradation regarding MMT mass fraction. A multidetector 390-LC apparatus from Varian with refractive index detectors was used with two 300 x 7.5 mm columns of PLgel 5 μm MIXED-D. Chloroform was used as a solvent for the samples, and as an eluent. Columns and detectors were held at 30 °C. Sample solutions were prepared to concentrations of approximately 2 mg/mL, and filtered with polytetrafluoroethylene filters with 0.45 μm pores. Weight average molar masses are measured regarding polystyrene standards. Tensile tests were conducted for each formulation, after 3D printing or after injection-moulding to compare the difference of tensile properties between additively manufactured materials and materials manufactured by a conventional formative process. 3D printed samples were tested at the ten configurations of the Doehlert network for mechanical optimisation. An Instron 33R 4204 apparatus was used with 1 kN load at a 5 mm/min constant tensile speed. Data were treated with Bluehill software. All measurements were repeated at least five times for mean response and standard deviations. Dixon tests (Qtests) were performed to eliminate aberrant values. Models associated to Doehlert designs were calculated after data processing and statistically checked by analysis of variance. Capillary rheology data were obtained after compounding to assess the effect of MMT mass fraction on the rheological behaviour thanks to a Rosand apparatus with two 1 mm capillary dies of 0.25 and 16 mm in length at 190 °C. Bagley and Rabinovitsch corrections were used to get viscosity as a function of the shear rate. Each formulation was measured at least two times, and mean values were obtained.

The microstructure was determined for each formulation and after each process to observe the developed MMT morphology regarding the formulation and the successive extrusions. A D8 Advance X-ray diffractometer was used at ambient temperature with a Cu Kα X-ray source of 0.15406 nm wavelength, generated at 40 kV and 40 mA. Measurements were performed with a 2Θ range between 2 and 10° with 2 s steps of 0.06°. Samples were in a powder form using a ZM1000 ultra-centrifugal mill under liquid nitrogen. TEM tests were conducted by the Institute of Materials Science of Mulhouse (IS2M, France) with a Philips CM200 at an acceleration voltage of 200 kV. TEM samples 90 nm thick were prepared using a Leica EM UC7 ultramicrotome at room temperature.

Results and discussion

3.1. Nanocomposite behaviour in processes • PLA molar mass stability Weight average molar mass of PLA in nanocomposite pellets after compounding is depicted in function of the MMT mass fraction in Figure 2. Considering experimental errors, a linear decrease with the mass fraction is observed, which was attributed to the presence of interlayer water. As the OMLS was not dried to increase the basal spacing and to improve the intercalation/exfoliation process [START_REF] Tenn | Effect of nanoclay hydration on barrier properties of PLA/montmorillonite based nanocomposites[END_REF], interlayer water could be released in the matrix during compounding and cause PLA to degrade if water content has reached 250 ppm [START_REF] Lim | Processing technologies for poly(lactic acid)[END_REF]. The decrease of the weight average molar mass could be essentially assigned to the degradation of PLA during compounding of the masterbatch since it was dried before dilution and compounding of nanocomposites from PLA05 to PLA7. This dilution could follow a proportionality rule and would explain the linear dependency. The decrease of PLA molar mass could significantly modify the rheological and thermomechanical properties of materials and should be reminded as an indirect effect of MMT on nanocomposite properties. Indeed, among others, direct effects of MMT are an increase of Young's modulus [START_REF] Zaidi | Relationship between structure and rheological, mechanical and thermal properties of polylactide/Cloisite 30B nanocomposites[END_REF], a higher thermal stability to degradation [START_REF] Carrasco | Kinetics of the thermal degradation of poly(lactic acid) obtained by reactive extrusion: Influence of the addition of montmorillonite nanoparticles[END_REF], and a thickening effect on polymer rheology [START_REF] Ray | New polylactide-layered silicate nanocomposites. 2. Concurrent improvements of material properties, biodegradability and melt rheology[END_REF], while a decrease of molar mass implies lower tensile properties, a lower stability to degradation [START_REF] Saiter | Probing the chain segment mobility at the interface of semi-crystalline polylactide/clay nanocomposites[END_REF] and a reduced viscosity [START_REF] Singh | Morphology and rheological behavior of polylactic acid/clay nanocomposites[END_REF]. • Evolution of MMT morphology Table 1 summarises XRD data and the intercalation state of MMT platelets in nanocomposites after each process step, namely compounding, filament shaping and printing by FFF (compound data from a previous work [START_REF] Ginoux | Decrease in non-linear viscosity of a polylactide nanocomposite with regard to the clay volume fraction[END_REF]). PLA5 presents a mixed morphology with the presence of peaks attributed to tactoids and intercalated platelets. Compared to the basal spacing of C30B, 1.87 nm, tactoids of PLA5 show a lower average basal spacing of 1.68 nm. This reduced basal spacing could be due to either an organo-modifier departure towards the PLA matrix or a beginning of its degradation leading to shorter hydrocarbonated chains and narrower interlayer spacing. By diluting the MMT loading, PLA3 only shows a peak of intercalated platelets, while XRD peaks for PLA05 and PLA1 were both not observable. The absence of peaks means either a fully exfoliated state or a low particle density to emit a signal in XRD analysis [START_REF] Singh | Morphology and rheological behavior of polylactic acid/clay nanocomposites[END_REF]. To verify the exfoliation state of PLA05 and PLA1, direct imaging is needed. Figure 3 illustrates the MMT morphology of PLA05 and PLA1 from TEM micrographs. It appears the two diluted nanocomposites, namely PLA05 and PLA1, consist of exfoliated platelets, except for rare intercalated platelets in PLA1 after compounding, totally absent after filament shaping and printing. Both formulations can be qualified as having reached a good exfoliation level. Therefore, the intercalation/exfoliation level increases as the MMT loading decreases. The observation is consistent with the work of Aubry et al. [START_REF] Aubry | Rheological investigation of the melt state elastic and yield properties of a polyamide-12 layered silicate nanocomposite[END_REF] reporting that particles tend to stay in an aggregate form after a maximum packing volume fraction, around 1 vol%, which returns an intercalation process less efficient with semi-diluted nanocomposites. Figure 4 shows the evolution of MMT morphology after each process step for PLA3 and PLA5. PLA3 essentially consists of intercalated platelets after compounding and exfoliated ones after filament shaping and printing. The evolution towards a more exfoliated state is consistent with the disappearance of peaks in XRD for PLA3 after filament shaping and printing (see Table 1). Similarly, PLA5 TEM micrographs after compounding (Figure 4b) indicates the presence of tactoids and some intercalated platelets and the presence of the three morphologies after filament shaping and printing (Figures 4d and4f), namely tactoids and intercalated and exfoliated platelets. These observations are consistent with the decrease of intensity of XRD peaks from tactoids in PLA5 after each process step and the appearance of a peak from intercalated platelets after printing (Table 1). Therefore, the successive extrusions tend to increase the intercalation/exfoliation state, which concurs with the works of Vaia and Giannelis [START_REF] Vaia | Lattice model of polymer melt intercalation in organically-modified layered silicates[END_REF] since the intercalation is a kinetic process. The increase of time spent in a molten state would then lead the nanocomposites to evolve towards a more intercalated or exfoliated state. However, a second XRD peak is observed for PLA5 around 7.6° after filament shaping and printing. The corresponding basal spacing is around 1.15-1.17 nm and close to that of native MMT (1.17 nm) [START_REF] Bordes | Nano-biocomposites: Biodegradable polyester/nanoclay systems[END_REF]. Consequently, PLA5 is shared between (1) an evolution towards a more intercalated/exfoliated morphology with the successive extrusions and lengthening of time spent in a molten state, and (2) the organo-modifier departure from the interlayer spacing or its degradation leading to more aggregated particles. The first phenomenon would develop the thermomechanical properties of the materials, while the second one would limit them [START_REF] Paspali | The influence of nanostructure on the mechanical properties of 3D printed polylactide/nanoclay composites[END_REF]. More precisely, the trend to aggregate could be an issue during printing by clogging the nozzle at high MMT loadings. Nevertheless, the reduced interlayer spacing is not accompanied by an increase of the number of particles in a stack, and the phenomenon seems then limited on TEM micrographs (Figure 4f) compared to the evolution towards a more intercalated state. • Evolution of MMT loading Figure 5 illustrates the evolution of MMT loading in nanocomposites after each process step. The MMT mass fractions appear close to target values after compounding: 0.5 ± 0.1, 1.0 ± 0.3, 2.7 ± 0.2, and 4.8 ± 0.3 wt% for, respectively, PLA05, PLA1, PLA3, and PLA5. A mild lower value for PLA3 and PLA5 could be assigned to an MMT loss during compounding. A more visible decrease of around 10% is observable for PLA3 and PLA5 from beginning to end, mostly assigned to the filament shaping process; MMT mass fractions were 2.5 ± 0.1 wt% for PLA3 and 4.3 ± 0.1 wt% for PLA5. Regarding standard deviations, FFF process does not significantly decrease the MMT mass fractions, with measured values being 2.4 ± 0.2 wt% for PLA3 and 4.2 ± 0.1 wt% for PLA5. Therefore, one should overload the polymer matrix by around 10%, depending on the used shaping process, to compensate the loss of MMT during filament shaping, and eventually during compounding. As a second consequence of these observations, nanocomposites do not present a high clogging risk for long duration prints since FFF process does not seem to remove MMT from materials to machine. Then, viability of nanocomposites with FFF is not expected to be limited by an accumulation of MMT in the printing head. • Nanocomposite behaviour during melting inside printing head Figure 6 shows the DSC analysis that informs on the behaviour of nanocomposites during melting, which would determine the temperature of use. As the filaments after shaping were quenched, the polymer matrix was amorphous. The temperature of use of nanocomposites in FFF could be then just above the softening temperature, above the glass transition after 76 °C (see Figure 6). However, DSC measurements also indicate that nanocomposites go through a cold crystallisation, starting around 89 °C, which would increase the viscosity of materials and cause a greater probability of nozzle clogging. Cold crystallisation temperature of nanocomposites is lower than that of neat PLA and could be caused by a nucleating effect of the MMT [START_REF] Di | Poly(lactic acid)/organoclay nanocomposites: Thermal, rheological properties and foam processing[END_REF], which would ease the cold crystallisation mechanism. A decrease of the molar mass with the clay loading (see Figure 2) could also be an explanation to the ease of nanocomposite to crystallise because of the greater mobility of polymer chains [START_REF] Vasanthakumari | Crystallization kinetics of poly(l-lactic acid)[END_REF]. Around 160 °C, an exothermic peak is visible, which is attributed to a disorder-to-order transition [START_REF] Saiter | Probing the chain segment mobility at the interface of semi-crystalline polylactide/clay nanocomposites[END_REF][START_REF] Zhang | Disorder-to-order phase transition and multiple melting behavior of poly(Llactide) investigated by simultaneous measurements of WAXD and DSC[END_REF], and is followed by the melting of PLA around 180 °C. The variation in melting temperature between the different formulations is not significant, which is consistent with other PLA-based nanocomposites, namely PLA/graphene for instance [START_REF] Wu | Crystallization Behavior of Polylactide/Graphene Composites[END_REF]. However, greater melting temperatures were also reported with other nanocomposites, namely polyether block amide/graphene nanocomposite, while adding nanofiller [START_REF] Sarwar | Fibrous PEBA-graphene nanocomposite filaments and membranes fabricated by extrusion and additive manufacturing[END_REF]. From those two studies, such difference is explained by the nanofiller/matrix interaction and the degree of intercalation/exfoliation of the nanoplatelets [START_REF] Wu | Crystallization Behavior of Polylactide/Graphene Composites[END_REF][START_REF] Sarwar | Fibrous PEBA-graphene nanocomposite filaments and membranes fabricated by extrusion and additive manufacturing[END_REF]. In the case of PLA-based nanocomposites, when observing the morphology of PLA/OMLS (see Table 1 and Figure 3), the low number of platelets in diluted systems and the transition to an intercalated morphology in the semi-diluted systems are not sufficient to impact the melting process (see Figure 6). Consequently, the printing temperature should be above the melting temperature, which ends up after 190 °C independently of the formulation, to avoid a potential clogging risk because of the crystallisation. The temperature of use of all PLA/OMLS nanocomposite filaments is recommended to be above 190 °C. • Printability Printability is conditioned by the possibility of use of filaments and the rheological behaviour of nanocomposites in extrusion process. Firstly, as written in section 2. Materials and methods, PLA7 cannot be printed because of the filament stiffness increasing with the MMT loading, as already reported by Lee et al. [START_REF] Lee | Thermal and mechanical characteristics of poly(l-lactic acid) nanocomposite scaffold[END_REF]. Secondly, the increase of MMT mass fraction is expected to increase the viscosity [START_REF] Mueller | The rheology of suspensions of solid particles[END_REF], which would cause the impossibility of the materials to be printed [START_REF] Ivanova | Additive manufacturing (AM) and nanotechnology: promises and challenges[END_REF]. Nevertheless, semi-diluted nanocomposites were observed with a lower viscosity compared to that of neat polymer in nonlinear rheology [START_REF] Gu | Melt rheology of polylactide/montmorillonite nanocomposites[END_REF][START_REF] Cho | Nylon 6 nanocomposites by melt compounding[END_REF]. These authors suggested a possible slip between the matrix and exfoliated MMT as well as a decrease in molar mass with the addition of fillers during mixing. Despite the increase of viscosity with the addition of fillers, it appeared according to Ginoux et al. study [START_REF] Ginoux | Decrease in non-linear viscosity of a polylactide nanocomposite with regard to the clay volume fraction[END_REF] that the decrease in viscosity is mainly attributed to greater activation energies with semi-diluted nanocomposites and to the PLA molar mass degradation as observed in Figure 2. Therefore, the viscosity in function of the MMT loading increases after the first addition of MMT, then decreases up to be lower than that of neat PLA after 3 wt% of MMT (see Figure 7). The decrease is more pronounced with higher printing temperature since the dependency of viscosity on temperature is greater with semi-diluted nanocomposite (> 1 wt%) [START_REF] Ginoux | Decrease in non-linear viscosity of a polylactide nanocomposite with regard to the clay volume fraction[END_REF]. This observation is in agreement with the printing failure rate noted during experimental design. On the entire Doehlert network, consisting of ten experiments, PLA0 showed a global printing failure rate around 10%, while that of PLA05 and PLA1 was between 20 to 40% in function of the configuration. The higher failure rate is consistent with a higher probability of nozzle clogging because of the thickening effect by the filler with diluted systems. On the contrary, PLA3 and PLA5 showed a drastic decrease of printing failure rate, below 10%, which is, again, consistent with a thinning effect by the filler with semi-diluted systems. The printability is then improved for nanocomposites loaded at more than 3 wt% of MMT due to a direct and indirect thinning effect by the filler on the molten materials but worsened at more than 7 wt% due to the stiffness of the filaments.

Figure 7.

Rheogram plotting the viscosity against the shear rate for PLA0 to PLA5 at 190 °C. Black arrows show the increase followed by the decrease of viscosity with the increase of MMT loading.

3.2.

Thermomechanical optimisation • Thermal stability Figure 8 depicts the evolution of thermal degradation temperatures of PLA with MMT loading. The more the MMT is added to PLA, the lower the onset of degradation, and the greater the remaining residue. Actually, the addition of plateletlike fillers was reported in the literature to shift the onset of degradation towards higher temperatures by inducing fewer effective collisions to trigger the intra-and inter-molecular degradation phenomena [START_REF] Zaidi | Relationship between structure and rheological, mechanical and thermal properties of polylactide/Cloisite 30B nanocomposites[END_REF][START_REF] Carrasco | Kinetics of the thermal degradation of poly(lactic acid) obtained by reactive extrusion: Influence of the addition of montmorillonite nanoparticles[END_REF]. However, Saiter et al. [START_REF] Saiter | Probing the chain segment mobility at the interface of semi-crystalline polylactide/clay nanocomposites[END_REF] also noticed a decrease of the onset of degradation with the addition of MMT, which was supposed to be due to the predegradation of the polymer matrix by the addition of filler during compounding. As seen in Figure 2, the addition of OMLS to the PLA reduced the matrix molar mass and is consistent with the assumptions from Saiter et al. study [START_REF] Saiter | Probing the chain segment mobility at the interface of semi-crystalline polylactide/clay nanocomposites[END_REF]. The endset of degradation firstly indicates an increase with the addition of MMT from 0 to 3 wt%, then a decrease afterwards. The degradation temperature range appears more reliable for discussion as the endset is dependent on both the onset and the kinetics of thermal degradation. In that way, the degradation temperature range is around 38.6 ± 1.4 °C for diluted systems (from PLA0 to PLA1 included) and reaches 49.2 ± 0.5 °C for semi-diluted systems (PLA3 and PLA5). Krishnamachari et al. [START_REF] Krishnamachari | Biodegradable poly(lactic acid)/clay nanocomposites by melt intercalation: A study of morphological, thermal, and mechanical properties[END_REF] had suggested the migration of MMT to explain the increase of thermal stability since the filler would create a protective physical barrier. A direct visual observation after TGA of diluted systems showed the presence of a white powder with granulometry in the micrometer range, while semi-diluted ones had a different aspect (see Figure ). PLA3 and PLA5, after TGA, showed a remaining porous and friable structure of the size of the pellets before thermal degradation. This structure could act as a physical barrier. As observed in Figure 9b, blisters are visible on the external surface and could reflect the overpressure by the accumulation of degraded materials. This phenomenon would result in a slower kinetics of thermal degradation by slowing down the departure of degraded materials to the outside of the system. The formation of such structure could be attributed to the relatively high amount of MMT in semi-diluted systems, which had previously indicated an increase of the filler-filler interactions at the expense of filler-matrix ones with a less exfoliated morphology (see section 3.1.), leading to a percolation. Consequently, the thermal stability is indirectly reduced by the predegradation of the PLA matrix by the OMLS inherent to the compounding process but is improved by a barrier-structure during degradation after a percolation threshold. After this critical mass fraction, kinetics do not seem modified and an optimal MMT loading is expected around 3 wt% to increase the thermal stability by the barrier-structure formation without reducing it excessively by the pre-degradation.

(a) (b) • Optimisation of tensile properties by Doehlert designs Elongations at break and at ultimate tensile strength were treated by Doehlert method, but experimental standard deviations were high, leading to low significance on all the coefficients and models. Therefore, printing speed and extrusion temperature have been concluded with no impacting effect on elongations. However, by taking the mean value of elongation at break from all the Doehlert design configurations for a same nanocomposite, then by comparing them in function of the MMT loading, the effect of the formulation is observable. The variation of elongation at break in function of the MMT loading for printed samples is compared that of injected moulded samples in Figure 10. One could see that, independently of the process, diluted systems show greater elongations at break than that of neat PLA, while semi-diluted ones show the lowest elongations. The observation is consistent with literature: Lai et al. [START_REF] Lai | Unusual mechanical properties of melt-blended poly(lactic acid) (PLA)/clay nanocomposites[END_REF] reported an extended elongation at break for PLA charged with 1 wt% of MMT compared to more loaded nanocomposites. This phenomenon is assigned to a plasticised interphase surrounding nanoparticles. This interphase modifies the micromechanical properties when abundant, which would be the case when particles are smaller than 100 nm [START_REF] Odegard | Modeling of the Mechanical Properties of Nanoparticle/Polymer Composites[END_REF]. With the use of MMT as a filler, 1 mm 1 mm only exfoliated platelets correspond to particles smaller than 100 nm, and intercalated or aggregated ones would not develop abundant plasticised interphase. As observed previously, only PLA05 and PLA1 have indicated a fully exfoliated state and would explain the extended elongations at break. For semi-diluted systems (PLA3 and PLA5), MMT stiffens the PLA-based materials and the elongation at break is reduced. However, the decrease is less pronounced for printed specimens than for injected ones that have half the elongations of neat PLA (Figure 10). The MMT loss during filament printing and the more exfoliated/intercalated morphology with the successive extrusions could be an explanation but did not seem to evolve to the point of increasing the elongation at break of the printed tensile test specimens from + 42% (PLA5) to + 97% (PLA3) compared to the injected samples (see section 3.1.). A more relevant explanation may come to the mesostructure inherent to the FFF process. The interlayer adhesion is known to be weak in additively manufactured parts, which brings brittleness to the parts [START_REF] Davis | Mechanical strength of welding zones produced by polymer extrusion additive manufacturing[END_REF]. This phenomenon is even more pronounced with the addition of filler as the nanocomposites have greater viscosities when shear rate is null, right after deposition, due to a percolation threshold [START_REF] Ginoux | Decrease in non-linear viscosity of a polylactide nanocomposite with regard to the clay volume fraction[END_REF], which causes a quite weak adhesion of the deposited filaments [START_REF] Bellehumeur | Modeling of Bond Formation Between Polymer Filaments in the Fused Deposition Modeling Process[END_REF]. The enhanced elongation at break is then counter-intuitive, but an explanation stands in the de-adhesion of the rasters in a same layer, rather than between layers [START_REF] Davis | Mechanical strength of welding zones produced by polymer extrusion additive manufacturing[END_REF]. As the materials are deposited in a 45°/-45° fashion and due to the structure consisting of deposited filaments, the rasters separate from each other by shearing, which realigns each raster in the load direction during tensile testing and the tensile test specimens would be able to fracture filament by filament (see Figure 11). The fracture toughness is then improved for printed samples, while injected ones have a sharper fracture for the semidiluted systems, namely PLA3 and PLA5. Therefore, the mesostructure of the printed samples would compensate the loss of elongation at break caused by the addition of a high MMT amount in PLA with the injected moulded specimens. Young's modulus and ultimate tensile strength were chosen as responses for models with Doehlert designs. Models globally showed a minimum or a tendency to a minimum on the first design (15 to 75 mm/s and 190 to 240 °C) and a maximum or a tendency to a maximum on the second design (45 to 105 mm/s and 190 to 240 °C) as already reported in Ginoux et al. work [START_REF] Ginoux | Influence of fused filament fabrication parameters on tensile properties of polylactide/layered silicate nanocomposite using response surface methodology[END_REF] because of the numerous phenomena modifying tensile properties. To compensate the loss of mechanical properties from formative to additive process, maximum responses are desired for optimisation and can be calculated from Doehlert quadratic mathematical model. Predicted maximum responses and factor values corresponding to the maximum response are summarised in Table 2. Printing speed and extrusion temperature tend to increase then to decrease with the addition of OMLS in PLA (see Table 2). The variation is analogous to that of viscosity (see Figure 7). Because of several deposition defects depending on flowability and instabilities with the printing speed, an optimal viscosity is expected [START_REF] Ginoux | Influence of fused filament fabrication parameters on tensile properties of polylactide/layered silicate nanocomposite using response surface methodology[END_REF], and is an explanation to the increase then to the decrease of printing speed, correlated to the shear rate. As nanocomposites show shear-thinning behaviour (see Figure 7), higher shear rates compensate the higher viscosities with diluted systems (PLA05 and PLA1) compared to semi-diluted systems (PLA3 and PLA5) or neat PLA. One should note that extrusion temperature tends to increase with printing speed. An explanation stands in the decrease of residence time in the heating nozzle with higher flow rates, which is not favourable to an optimal heat transfer. Although higher shear rates could self-heat the molten materials in opposition, this phenomenon does not compensate the lack of heat transfer and would need a compensation by a higher extrusion temperature setpoint. 12 compares the Young's modulus and the ultimate tensile strength between full and printed specimens at different MMT loadings. As expected, according to literature [START_REF] Ngo | Additive manufacturing (3D printing): A review of materials, methods, applications and challenges[END_REF], tensile properties are lower for printed samples than for injected ones. The properties of the printed specimens are represented with the maximum response and the minimum response obtained throughout the Doehlert network. These responses are experimental and not predicted. These values may then extend beyond (see Table 2). The average mechanical properties over the entire network are also displayed in Figure 12 and serve as a basis for evaluating the variation of the properties with the MMT loading, independently of a specific configuration. Whatever the manufacture process, the Young's modulus is linearly correlated with the MMT loading (Figure 12a). However, the improvement of the Young's modulus with the MMT loading is enhanced for the injected specimens with a sharper slope compared to that of the printed specimens. Conversely, the maximum stress decreases linearly with the MMT loading for the printed specimens (Figure 12b), and the use of nanocomposite cannot compensate the loss of properties from injected samples to printed ones. Similarly to elongation at break, diluted systems show a synergetic effect with the development of matrix/filler interaction over filler/filler or matrix/matrix interactions, which lead to higher ultimate tensile strength compared to ultimate tensile strengths for neat PLA or semi-diluted systems. This synergetic effect is not observable with printed samples and could be hidden by the dominance of structural defects inherent to the FFF process. For neat PLA, the average loss of mechanical properties between formative and additive processes is estimated around 14.6% for the Young's modulus and 11.3% for the ultimate tensile strength. However, this loss can reach respectively 19.0% and 13.6%, but can also be minimised to 10.3% and 7.4% regarding the configuration in the studied Doehlert network. With the prediction from the theoretical mathematical models (see Table 2), this loss of properties can be lowered to 9.9% and 7.2% for Young's modulus and ultimate tensile strength. Consequently, the optimisation of the printing speed and extrusion temperature cannot reach alone the mechanical properties of the injected PLA0 specimens. Nevertheless, the use of the Doehlert method divides by two the loss of mechanical properties by comparing an unfavourable configuration with a favourable one. Full property compensation must then go through the use of nanocomposite. Thus, the Young's modulus of injected moulded PLA0 is approximately equal to that of PLA3 at its maximum response and that of PLA5 at its minimum response. On the other hand, because of the negative tendency of the ultimate tensile strength with the MMT loading, the minimum loss of property is achieved with the PLA0. To obtain similar properties to the PLA0 specimens manufactured by injection moulding, one should opt for the more diluted nanocomposite in MMT to avoid the loss on the ultimate tensile strength while reaching the same Young's modulus at a configuration maximising the response. Therefore, the PLA3 printed to the maximised configuration seems to be the best compromise if the optimisation of the Young's modulus and/or thermal stability (see Thermal stability section) has priority over the ultimate tensile strength. In addition, its elongation at break is close to that of injected PLA0. For a priority optimisation on the ultimate tensile strength, the PLA0 printed to the maximised configuration is recommended.

Conclusions

Characterisation results were compared with those resulting from the Doehlert experimental design methodology and revealed (i) an optimisation of thermal stability and tensile mechanical properties via the parametric printing process depending on the formulation, and (ii) lower viscosities due to a pre-degradation of the polymer matrix with the addition of OMLS during compounding and to a greater dependency of viscosity on temperature with semi-diluted systems, accompanied by a lower nozzle clogging probability and so a lower failure rate and a better printability for PLA charged at intermediate loadings. The presented methodology allows a deeper mastery of the manufacturing chain by simultaneously controlling materials formulation and process-induced structure to freely tailor thermomechanical performance of additively manufactured composites. Numerical modelling based on this experimental method would help design parts by predicting multi-physics characteristics, which would reduce the number of test-failure iterations in applications where waste must be limited, such as, for example, in space station with on-demand tooling or repairing. Several recommendations are presented in this work. Despite an improved fracture toughness with printed nanocomposites compared to injected moulded nanocomposites, future studies should focus on new formulations with plasticiser since PLA remains a brittle matrix. As another way, formulations with nanohydroxyapatite could help design functional scaffolds for bone tissue engineering applications in combination with barrier properties of MMT for drug-controlled release. Future studies should also focus on the filament ageing. As MMT is a nucleating agent and the filaments are amorphous, a crystallinity could appear on the long term and cause the filament to gain brittleness. Other additives should be sought to enhance the viability of functional nanocomposites with FFF.

Figure 1 .

 1 Figure 1. Schematic diagram of the overall experimental section, from materials preparation to characterisation tests.

Figure 2 .

 2 Figure 2. Linear decay of the weight average molar mass of PLA, as measured by GPC, with the addition of OMLS during the elaboration of nanocomposites. Horizontal bars depict the standard deviations of MMT mass fractions as determined by TGA. Vertical bars are an assumed error of 10% on molar mass measurements.

  mol)MMT mass fraction (wt%)

Figure 3 .

 3 Figure 3. TEM micrographs of (a) PLA05 and (b) PLA1 nanocomposites after compounding. Black oval points out some observed intercalated platelets in PLA1.

Figure 4 .

 4 Figure 4. TEM micrographs of (a,c,e) PLA3 and (b,d,f) PLA5 nanocomposites after (a,b) compounding, (c,d) filament shaping, and (e,f) printing.

Figure 5 .

 5 Figure 5. Evolution of MMT mass fraction after compounding, filament shaping and printing for each formulation. Vertical bars represent the standard deviations calculated from MMT mass fraction measurements as determined by TGA.

Figure 6 .

 6 Figure 6. DSC thermograms of PLA0 (red), PLA05 (green), PLA1 (blue), PLA3 (black), and PLA5 (purple) after filament shaping at the first heating.

Figure 8 .

 8 Figure 8. TGA thermograms at a 10 K/min heating rate from PLA0 to PLA5. Black arrows show the variation of thermal degradation onset and of residual mass after degradation with the increase of MMT loading in PLA.

10

 10

Figure 9 .

 9 Figure 9. Photographies of a PLA5 pellet after thermal degradation, cut in half with (a) the inside view and (b) the outside view.

Figure 10 .

 10 Figure 10. Comparison of the elongations at break of 5A-type tensile test specimens manufactured by injection moulding (in light orange) and by FFF (in dark brown) in function of the MMT loading. Vertical bars represent the standard deviations on ten measures for injected moulded specimens and on fifty measures for printed specimens (five measures at each of the ten configurations of the Doehlert network).

Figure 11 .

 11 Figure 11. Photography of the fracture of a printed 5A-type tensile test specimen of PLA3 after tensile test at a 5 mm/min tensile speed.

Figure 12 .

 12 Figure 12.Comparison of (a) the Young's moduli and (b) the ultimate tensile strengths in function of the MMT mass fraction in PLA, as measured by tensile test on injected moulded 5A-type specimens (blue diamond) and on printed 5Atype specimens at the maximum response (green circle), at the minimum response (red square) and at the mean response (orange triangle) of the ten configurations from the Doehlert network. Vertical bars represent the standard deviations on ten measures for injected moulded specimens and on five measures for printed specimens.

Table 1 .

 1 Comparisons of Bragg angle, intensity, and basal spacing of MMT platelets in C30B and nanocomposites at different MMT contents from 0.5 to 5 wt% and between each process step, as determined by XRD measurements.

		Sample name	2Θ (°)	Intensity (cps)	𝒅 𝟎𝟎𝟏 (nm)	Comments
	Control	C30B	4.73	18207	1.87	Tactoid
		PLA05	/	/	/	/
		PLA1	/	/	/	/
	Compound	PLA3 2.89	323	3.05 Intercalated
		PLA5 2.59	153	3.41 Intercalated
		PLA5 5.26	698	1.68	Tactoid
		PLA05	/	/	/	/
		PLA1	/	/	/	/
	Filament	PLA3	/	/	/	/
		PLA5 5.03	499	1.76	Tactoid
		PLA5 7.52	170	1.17	Tactoid
		PLA05	/	/	/	/
		PLA1	/	/	/	/
	Print	PLA3 PLA5 2.59 /	/ 81	/ 3.41 Intercalated /
		PLA5 5.20	300	1.70	Tactoid
		PLA5 7.69	74	1.15	Tactoid

Table 2 .

 2 Maximum responses predicted by Doehlert design with the Young's modulus (E) and the ultimate tensile strength (σM) as response in function of the MMT loading and with the printing speed and extrusion temperature associated to the predicted response. Prediction is not presented if the maximum is out of the Doehlert network.

	Sample	Property	Speed (mm/s)	Temp. (°C)	Prediction (MPa)
	PLA0	E σM	67 73	212 209	1274 64,0
	PLA05	E σM	~ 60 60	> 240 221	/ 62,2
	PLA1	E σM	95 ~ 105	239 > 240	1268 /
	PLA3	E σM	81 67	217 206	1401 59,6
	PLA5	E σM	72 ~ 50	192 < 190	1527 /
	Figure

Acknowledgements

The project PolyFabAdd is co-funded by the European Union. Europe invests in Champagne-Ardenne with the European Regional Development Fund. The authors gratefully acknowledge the Région Champagne-Ardenne (France) (grant number D201501201) and the European Union for their financial support (grant number D201505025). This project is also included in the MATUR Chair, co-financed by the Région Champagne-Ardenne (France) (grant number D201507798), the European Union (grant number D201507799) and UIMM (Partners convention n°13-2015). Dr. Florica Lazar and Mr. Stéphane Buet are thanked for their technical support in this project.


Pierre Gras

Christian Rey

Olivier Marsan

Stéphanie Sarda

Christèle Combes

Synthesis and Characterisation of Hydrated Calcium Pyrophosphate Phases of Biological Interest

Keywords: Biological calcification, Physicochemical characterization, Calcium pyrophosphate, Synthetic methods, Hydrates, Vibrational spectroscopy

The details of a synthesis method for biologically relevant hydrated calcium pyrophosphates (CPPs, Ca 2 P 2 O 7 •nH 2 O) has been elucidated. Control of the pH (from 3.6 to 5.8) and the temperature (from 25 to 90 °C) during the synthesis enabled the preparation of four pure CPP phases within one hour without intermediates: monoclinic and triclinic calcium pyrophosphate dihydrate (CPPD, Ca 2 P 2 O 7 •2H 2 O), which are the two CPP phases detected in vivo in joints of arthritic patients, monoclinic tetrahydrate β (CPPT, Ca 2 P 2 O 7 •4H 2 O) and an amorphous phase (a-CPP, Ca 2 P 2 O 7 •nH 2 O). Four domains corresponding to the four different phases of hydrated calcium pyrophosphate were identified; a-CPP was synthesised over a very wide pH and temperature range (up to 90 °C) within the domain of synthesis conditions explored, including physiological conditions (pH 7.4 and 37 °C). The as-synthesised hydrated CPP phases were characterised by complementary techniques (powder X-ray diffraction, FTIR and Raman spectroscopy, scanning electron microscopy and thermogravimetry) and chemical analyses. Rietveld refinement analyses of the as-synthesised crystalline phases were performed, and

Introduction

Calcium pyrophosphate hydrates (CPP, Ca 2 P 2 O 7 •nH 2 O) are involved in several forms of arthritis, including calcium pyrophosphate crystal deposition disease, also known as pseudogout, and osteoarthritis (OA), the most frequent form of rheumatic disease. [1,2] To date, two different polymorphs of dihydrate CPP crystals are commonly identified in joint tissues of arthritic patients: monoclinic and triclinic calcium pyrophosphate dihydrate (CPPD, Ca 2 P 2 O 7 •2H 2 O) crystals, referred to as m-CPPD and t-CPPD, respectively. They have been reported in a large range of sizes, typically from 0.06 to 0.3 μm in the temporomandibular joint and from ca. 1 to 20 μm in the knee. [3,4] They both appear to induce an inflammatory response, possibly because of the rupture of lysosome phospholipid membranes induced by there were significant differences between the m-CPPD Xray diffraction pattern observed and previously published cell parameters. Vibrational spectroscopy allowed the crystalline and amorphous phases synthesised to be clearly distinguished and identified owing to the high flexibility of the pyrophosphate anion. Chemical analyses showed that the synthesis conditions used in this study did not allow significant hydrolysis of the pyrophosphate ions into phosphate ions, and the number of water molecules associated with each synthesised CPP phase was determined by thermogravimetric analysis. Different mechanisms of dehydration were also identified. The study of the formation of synthetic and well-characterised hydrated calcium pyrophosphate phases and their availability in large amounts in vitro could allow progress to be made on the biological role of these phases and their possible transformations. This could also aid their detection in patients suffering from disease caused by calcium salt crystals and could clarify the mechanism by which CPP crystals form and evolve in vivo.

pyrophosphate groups on the surface of the crystals. [5][6][7][8] However, current treatments cannot inhibit the formation of CPP crystals, although they help to resolve the inflammatory responses. [9][10][11] Although many papers have focused on CPP crystals associated with OA from a clinical point of view, very little in the way of physicochemical data on these phases of biological interest has been published. Several in vitro syntheses of hydrated CPP have been reported. [12][13][14][15][16][17][18][19][20] They were mainly based on the method developed by Brown et al. [12] and its modifications by Mandel et al. [13] Four calcium pyrophosphate hydrates were reported among the 25 calciumcontaining pyrophosphate compounds described in the literature, [12] including m-CPPD, t-CPPD, a dimorphic monoclinic tetrahydrate (CPPT, Ca 2 P 2 O 7 •4H 2 O) referred to as m-CPPT α and m-CPPT β. [21,22] However, these methods present some drawbacks as they lead to low total reaction yields and generally involve unstable calcium pyrophosphate intermediates, as highlighted by Groves et al. [20] Other methods were developed to overcome either or both of these disadvantages, especially for the synthesis of t-CPPD, m-CPPT β and an amorphous phase (a-CPP), but not for m-CPPD. [20,23,24] However, the crystalline phases are obtained with synthesis durations between two hours and one week, and the reproducibility of some of these methods is rather unpredictable and the pure phases are not always obtained. [25,26] There is an interest in studying the conditions of formation of all these hydrated CPP phases in vitro and in finely characterising them to better understand their formation in vivo (direct formation or possibly the hydrolysis and evolution of a metastable precursor phase). We recently reported the synthesis of m-CPPT β at pH 4.6 and 25 °C by a fast and simple synthesis method. [27] On the basis of this synthesis method, we could suggest that a variation of pH and temperature leads to the formation of pure pathological CPP phases, that is m-CPPD and t-CPPD.

In the present study, we investigated the formation conditions of hydrated calcium pyrophosphate phases obtained by a single-step precipitation in aqueous media. The four different forms obtained, m-CPPD, t-CPPD, m-CPPT β and a-CPP, were characterised by physicochemical methods and techniques including powder X-ray diffraction (XRD), FTIR and Raman spectroscopy, thermogravimetric and differential thermal analyses (TGA-DTA), scanning electron microscopy (SEM), the Brunauer-Emmett-Teller (BET) method for specific surface area determination and chemical analyses (calcium and phosphate titration).

Results and Discussion

Synthesis

The synthesis method implemented in the present study could be described as a direct (without intermediates) synthesis method for several hydrated calcium pyrophosphate phases. The double decomposition of the calcium nitrate and potassium pyrophosphate solutions occurs in an aqueous buffered medium containing ammonium and acetate ions. The precipitation reaction involved can be represented by Equation (1).

2Ca 2+ + P 2 O 7 4-Ǟ Ca 2 P 2 O 7 •nH 2 O (1)
Generally, the two most important parameters involved in the precipitation of various salts in different hydration states are temperature and pH. Our study focuses on the influence of these parameters, which were not systematically explored and described in previous works related to CPP synthesis. Although other parameters, such as stirring or reagent concentration, may determine the nature of the phases formed, most of our syntheses were performed with a stoichiometric calcium/pyrophosphate ratio of 2 and with magnetic stirring. Briefly, we investigated the influence of the method of addition of reagents into the buffer, especially the influence of concentration of the reagent solutions.

The influence of pH and temperature was studied by performing several experiments with pH varied from 3.6 to 5.8 and temperature varied from 25 to 90 °C at fixed concentrations. Samples from these experiments were analysed by FTIR spectroscopy and powder XRD (data not shown) and compared with data already published. [12,13,22,24,28] Figure 1 shows the formation domains of pure hydrated CPP phases for our synthesis method. At fixed concentrations of pyrophosphate and calcium (75 and 150 mmol L -1 , respectively, in reagent solution), we determined four domains corresponding to four different phases of hydrated calcium pyrophosphate: a-CPP, an amorphous phase, m-CPPT β, a tetrahydrated phase, both formed at low temperatures, and m-CPPD and t-CPPD, two biologically relevant phases synthesised at higher temperatures (Figure 1). Mixtures of these phases (generally two CPP phases, including a-CPP and poorly crystalline phases for pH above 4.5 and temperature under 70 °C) were obtained under pH and temperature conditions corresponding to the areas between the domains of pure CPP formation. The CPP crystalline compounds synthesised under these conditions have coherent domains over 50 nm in at least one direction, as determined by using the Scherrer equation. No morphological differences have been observed with modification of temperature for each formation domain, but synthesis under more acidic conditions promotes growth of well-faceted plates for m-CPPT β and thicker and longer needles for t-CPPD.

Several syntheses of pyrophosphate hydrates with the general formula M 2 P 2 O 7 •nH 2 O, in which M is a bivalent metal, have been reported: these include condensation of hydrogen phosphate ions HPO 4 2-, [29][30][31][32] precipitation processes, the use of ion-exchange resins [START_REF] Brouzi | Phosphorus Sulfur Silicon[END_REF]34] and double decomposition. In these cases, the product was then allowed to stand for several hours to crystallise into the desired phase. Double decomposition has been used to form amorphous and then crystalline pyrophosphate compounds of manganese pyrophosphate, Mn 2 P 2 O 7 •nH 2 O, [35] amorphous cobalt pyrophosphate, Co 2 P 2 O 7 •nH 2 O, which was used as an intermediate for the synthesis of crystalline compounds such as Co 2 P 2 O 7 •6H 2 O, [36] and calcium pyrophosphate (CPP), which was formed as an amorphous hydrated phase a-CPP and then crystallised into two dimorphs of Ca 2 P 2 O 7 •2H 2 O (t-CPPD and m-CPPD) and two dimorphs of Ca 2 P 2 O 7 •4H 2 O (m-CPPT α and m-CPPT β). [12,20,23,24] The coprecipitation of calcium and pyrophosphate ions at room temperature and without monitored pH leads to an amorphous phase. [12,20,24] We showed that the amorphous phase can also be synthesised by our method within a wide range of temperatures and pH (from ambient temperature to 90 °C and from pH 5.8 to 7.4) with no evidence of crystallisation (Figure 1).

Interestingly, this amorphous calcium salt can precipitate at much higher temperature than amorphous calcium orthophosphate or amorphous calcium carbonate, which are also amorphous calcium salts of biological interest. Unlike these latter amorphous phases, which may crystallise readily even at room temperature, a-CPP did not crystallise during the 45 min of synthesis at 90 °C or if kept in the mother liquor for over 6 months at room temperature. This stability could be explained by the flexible structure and probably the size of the P 2 O 7 4-anion, which could enhance inhibition of crystallisation of a-CPP compared to the CO 3 2-or PO 4 3-anions in calcium carbonate and calcium phosphate salts, respectively. [24] The highly flexible P-O-P bond has been reported with angles varying from 123°to 180°in the solid state. [37] However, it has been reported that the amorphous phase then crystallised within one day to one week under specific conditions into the different crystalline forms already mentioned. [12,13,20] Thereby, the published synthesis methods of crystalline CPP are mostly based on the dissolution of an intermediate, amorphous CPP (a-CPP) or calcium dihydrogen pyrophosphate (CaH 2 P 2 O 7 ), followed by crystallisation into the CPP phase by using different synthesis conditions (pH temperature, concentration, agitation, time). These syntheses were performed for more than 24 h and led to low total reaction yields. [20] The method we set up for the m-CPPT β and t-CPPD syntheses used conditions already described to lead to this phase from a-CPP as intermediate: temperature below 50 °C and pH above 5 for m-CPPT β, and in acidic medium and the use of a steam bath for t-CPPD. [12,13,20,23] However, the synthesis described in this study was achieved in one step within 45 min, as with the other three phases synthesised, and was faster than the 24 h to one week needed for the previously published methods.

The synthesis method also allowed synthesis of the m-CPPD phase, a biologically relevant phase that has been observed in several joints of arthritic patients in addition to t-CPPD [4] and appears to have greater inflammatory properties than t-CPPD. [5,6] m-CPPD has been poorly studied in the literature; the synthesis methods already published often involve Mg 2+ ions or lead to a mixture of m-CPPD and t-CPPD phases. [12][13][14][15]25] Figure 1 shows that the conditions for m-CPPD synthesis were very narrow: lower pH values and temperatures during the synthesis lead to the formation of m-CPPT β instead of the dihydrate and higher pH values and temperatures lead to t-CPPD, as already mentioned. [12] The synthesis conditions of pH 5.8 and 90 °C allowed the formation of the m-CPPD phase, and the reaction time was fast enough to avoid the transition into t-CPPD mentioned by Brown et al. [12] In water, m-CPPD and m-CPPT β were reported to slowly evolve to the t-CPPD phase. However, we determined that if m-CPPD was kept in the mother liquor at 50 °C for over 1 month, the as-synthesised m-CPPD phase does not evolve into t-CPPD.

Interestingly, the synthesis of m-CPPD described in the present study occurred without the addition of Mg 2+ ions to the solution; however, the nature of the crystals formed was reported to be highly dependent on the concentration of Mg 2+ ions, which are the main impurity in commercial calcium salts. The calcium nitrate salt we used contained less than 100 ppm Mg 2+ ions (manufacturer analysis). High concentrations of magnesium ions are known to favour m-CPPD formation, as reported by Cheng et al., [15,38] and magnesium ions are used for this purpose in the method developed by Mandel et al. [13] However, pyrophosphate ions are approximately five times more effective in sequestering magnesium ions than calcium ions and could lead to impurities in the crystals. [39] This point could partially explain why it is difficult to obtain a pure m-CPPD phase and why the structure of this phase is still unresolved 50 years after its first preparation and XRD characterisation. [4] Thus, we briefly investigated the influence of different parameters on the synthesis of m-CPPD.

Different molar ratios of calcium/pyrophosphate (from

[Ca 2+ ] = [P 2 O 7 4- ] to [Ca 2+ ] = 4[P 2 O 7
4-] with [Ca 2+ ] = 150 mmol L -1 ) were used at fixed pH and temperature corresponding to the conditions for the formation of m-CPPD (pH 5.8 and 90 °C).

A mixture of m-CPPD and t-CPPD appeared at [Ca 2+ ] = 1.5[P 2 O 7

4-] and below, and an amorphous phase formed for [Ca 2+ ] = 3[P 2 O 7

4-] and above. It should also be noted that if the calcium solution was added to the buffered pyrophosphate reagent with [P 2 O 7 4-] = 75 mmol L -1 , that is, with a large excess of pyrophosphate ions compared to calcium ions, a phase identified as a pentacalcium diammonium pyrophosphate hexahydrate, Ca 5 (NH 4 ) 2 (P 2 O 7 ) 3 •6H 2 O, formed.

Physicochemical Characterisation of as-Synthesised Pure CCP Phases

Four samples, one for each pure CPP phase synthesised at pH and temperature within their respective domain of formation, were selected for further characterisation. These samples correspond to the synthesis conditions reported in Table 1. The crystalline samples selected were identified by powder X-ray diffraction analysis (Figure 2) as m-CPPT β, m-CPPD and t-CPPD by the characteristic pattern of each phase according to the published data (Table 2). [13,22,28] A fourth pattern, including a broad diffuse halo was also observed and is characteristic of an amorphous phase (a-CPP). The pattern obtained for the m-CPPD phase was in good agreement with the data published by Brown et al. and the patterns published by Liu et al. [12,19] but did not correspond to the lattice constant determined by Mandel et al. [13] The structure of m-CPPD has not been resolved, even though m-CPPD has been identified in arthritic joints. [5,25] Rietveld refinements were performed on the X-ray diffraction patterns of the as-synthesised m-CPPD and were based on the dihydrate pyrophosphate compound cell parameters from other monoclinic pyrophosphate dihydrate structures fully identified as Mg 2 P 2 O 7 •2H 2 O, [30] Co 2 P 2 O 7 • 2H 2 O, [31] Fe 2 P 2 O 7 •2H 2 O [32] and Mn 2 P 2 O 7 •2H 2 O, [35] which share very similar cell parameters and atomic organisation.

However, no satisfactory cell parameters similar to those of the X 2 P 2 O 7 •2H 2 O structures were found from the powder X-ray diffraction pattern; therefore, the m-CPPD structure could be significantly different from those of the other dihydrate pyrophosphate compounds (Table 2).

Table 2. Unit-cell parameters obtained by Rietveld refinement for the m-CPPT β, t-CPPD and m-CPPD samples compared to the single crystal data from Balic ´-Z ˇunic ´et al. and Mandel. [22,28] Estimated standard deviations are given in parentheses. 3. 4-for t-CPPD, [40] these vibrations could not be separated in the spectra, and the data are discussed with regard to the vibrational domains of the free ion. Eight P-O stretching modes can be distinguished in the following order of frequencies and assignments: four ν as PO 4 Ͼ ν s PO 4 Ͼ ν as PO 4 Ͼ ν as POP Ͼ ν s POP (Figures 3, a and4). P-O and P-O-P stretching vibration bands were observed between 1250 and 950 cm -1 and between 950 and 700 cm -1 , respectively, in the spectra of t-CPPD, m-CPPT β and m-CPPD (Table 3).

Ref

. R p /R wp a [Å] b [Å] c [Å] α [°] β [°] γ [°] m-CPPT
These P 2 O 7 4-vibration modes were observed for all three crystalline samples (t-CPPD, m-CPPT β and m-CPPD) by IR and/or Raman spectroscopy. The band positions are somewhat distinct and allowed unambiguous attribution of the bands for each hydrated CPP phase.

As already noted, the 21 predicted active modes are not all apparent, although several bands present shoulders. In particular, the ν as POP vibration in the FTIR spectra seems to be coupled with another band, and this is especially apparent for m-CPPT β. This band was not observed in the Raman spectra.

Water O-H vibrations occurred in the 3700-2800 cm -1 region, and there are significant differences between the spectra of the synthesised pure CPP phases. This region of the m-CPPD spectrum seemed closer to the spectrum of m-CPPT β, a tetrahydrate with layers composed of water molecules, than to that of t-CPPD, another dihydrate. A weak broad adsorption band also appeared in the 2400-1950 cm -1 range. In the absence of H 3 O + ions and P-O-H groups in the crystal, this band was presumed to correspond to composite vibration of water molecules (δ + ω) OH . [41] The spectroscopic data also allow direct information about the P 2 O 7 4-configuration to be obtained, and especially the P-O-P bond angle, as proposed by Rulmont et al. [37] In t-CPPD and m-CPPT β, the P 2 O 7 4-anions have a bent P-O-P angle and an eclipsed configuration. In particular, the P-O-P angle has been reported as 123.1°for t-CPPD and 134.1°for m-CPPT β, both in the dichromate configuration. [22,28] The P 2 O 7 4-anion could be considered as a three-body system, mainly by disregarding couplings between the vibrations of the P-O-P bridge and those of external PO 3 groups and the influence of Ca on the vibrational frequencies of the anion. Within this approximation, the stretching frequencies observed by IR and Raman spectroscopy depend on the P-O-P angle θ and the force constant k of the P-O bond.

Lazarev suggested an empirical relationship between the P-O-P angle θ and a quantity Δ [Equation ( 2)] that eliminates the contribution of the force constant k. [START_REF] Lazarev | Vibrational Spectra and Structure of Silicates[END_REF] Δ = (ν as POP -ν s POP)/(ν as POP + ν s POP) = f(θ)

(2)

The Lazarev relationship has been applied to pyrosilicates, pyrogermanates and pyrophosphates and was satisfactorily obeyed by a large majority of compounds, al-though a few of them exhibited a significant deviation from the main curve. [37] In those cases, the two approximations mentioned above could explain the differences observed. However, the relationship seems to be confirmed and is linear for pyrophosphate compounds with dichromate structure.

The Lazarev relationship was applied to m-CPPD with 100Δ = 11.54, and a θ angle of 133.6°was determined from the plot published by Rulmont et al. [37] This value was compared to those of the other calcium pyrophosphate compounds α-Ca 2 P 2 O 7 , β-Ca 2 P 2 O 7 , m-CPPT β and t-CPPD, which are 130.0, 130. 5 O (127.5, 125.9, 126.2 and 125.7°, respectively), all of which exhibited much lower values for the P-O-P angle. [30][31][32]35] Differences between pyrophosphate phase structures have already been reported for anhydrous compounds by Brown and Calvo. [START_REF] Brown | [END_REF] Crystals of the composition X 2 Y 2 O 7 (Y = Si, P, As, Ge) with an ionic radius for X of less than 0.97 Å (X = Mg, Mn, Fe, Co, Ni, Cu, Zn) are isostructural with thortveitite (Sc 2 Si 2 O 7 ) with X-O-X bond angles of 140-180°, an O-Y•••Y-O torsion angle of 60°(staggered conformation), and the YO 4 tetrahedra shows a very low degree of distortion, However, those in which X has an ionic radius greater than 0.97 Å (X = Ca, Sr, Ba) usually crystallise with a dichromate structure with a Y-O-Y angle of ca. 120-130°, an O-Y•••Y-O torsion angle of 0-30°( eclipsed conformation) and distorted YO 4 tetrahedra. [21] Unlike the corresponding anhydrous compounds, Mn 2 P 2 O 7 Moreover, Mandel correlated this value of the θ angle to a deformed tetrahedra and a P-O distance in the bridge between 1.61 and 1.63 Å. [28] According to Cruickshank, this relationship is based on the dπ-pπ overlap in the P-O bonds. [44] A P-O-P angle of ca. 130°would degrade dπ-pπ overlap in the P-O bonds in the bridge and, therefore, result in a larger P-O distance.

For an angle of ca. 180°, the P-O bonds receive some π character from the p orbital in the P-O-P plane as well as from the p orbital perpendicular to the plane, and the bond length then decreases. At the P-O-P angle limit of 180°, both p orbitals on the oxygen atom perpendicular to the P-O-P line would share fully in the two π systems. This case corresponds to the thortveitite Sc 2 Si 2 O 7 structure, in which the PO 4 tetrahedra show a very low degree of distortion, as in the β-Mg 2 P 2 O 7 structure.

For an angle of ca. 130°, the π character decreases and the bond length is supposed to increase. [45] A comparison of the Raman and IR bands of the PO 3 deformation modes between 700 and 400 cm -1 and the stretching modes between 1250 and 950 cm -1 could also be used to determine if the pyrophosphate compounds have a centrosymmetric structure. The coincidence of the Raman and IR bands in this domain confirms that the structure is not a centrosymmetric thortveitite structure.

Chemical analyses were performed to determine the composition of the as-synthesised CPP phases. We first considered that the conditions of synthesis used in these preparations of calcium pyrophosphate phases could lead to hydrolysis of pyrophosphate anions into two hydrogen phosphate anions [Equation ( 3)].

P 2 O 7 4-+ H 2 O h 2HPO 4 2- (3)
This hydrolysis reaction has been reported to occur through two different processes: as an intracrystalline process referred to as internal hydrolysis and a process in solution referred to as external hydrolysis. [24] Internal hydrolysis of pyrophosphate was observed in amorphous calcium pyrophosphate at a moderate temperature of 37 °C. Water molecules contained in the bulk solid react with nearby pyrophosphate molecules to form small amounts of phosphate anions. [24] However, the phenomenon is negligible under 140 °C.

At lower temperatures, a homogeneous hydrolysis of P 2 O 7 4-ions in solution occurs and is favoured by acidic pH. These reaction conditions were implemented for the chemical analysis of P 2 O 7 4-ions in the as-synthesised CPP samples (see Experimental Section). Significant amounts of orthophosphate anions could be produced by this process during the synthesis depending on the pH and temperature used. [46] Phosphorus titrations, of orthophosphate (PO 4 3-and HPO 4

2-

) and pyrophosphate (P 2 O 7 4-) ions, were performed to assess the percentage of pyrophosphate hydrolysed during the synthesis (Table 4). The results indicate that pyrophosphate hydrolysis is very limited: ca. 2 % of the P 2 O 7 4ions of each crystalline CPP phase synthesised were titrated as hydrolysed. This hydrolysis could have occurred during the preparation of the sample for titration or during the synthesis. In any case, the quantity of orthophosphate ions detected corresponds to the detection limit of the method. The present synthesis conditions of temperature and pH do not allow significant hydrolysis of pyrophosphate anions, even under the more extreme conditions (acidic pH and high temperature). The significantly higher concentration of orthophosphate ions in the a-CPP sample, synthesised at pH 5.8 and 25 °C, indicated either a propensity of the noncrystalline structure to incorporate phosphate ions or an easier hydrolysis of pyrophosphate ions in this compound after its formation.

The calcium to phosphorus ratio is commonly used to partially identify phases for calcium phosphate compounds and quantify calcium deficiency. Calcium titrations were performed and the results exhibited a characteristic Ca/P ratio of 1 for all calcium pyrophosphate compounds synthesised (Table 4). This indicates that compounds show no calcium deficiency, which could have happened, for example, if protonated pyrophosphate species were present in the precipitates. Moreover, the titration of the crystalline compounds was consistent with the expected water content in the crystalline phases: two H 2 O molecules for m-CPPD and t-CPPD and four H 2 O molecules for m-CPPT β.

Thermogravimetric analyses were performed to study the evolution of the calcium pyrophosphate hydrates up to 500 °C (Figure 5). The weight loss in the crystalline phases occurred in three steps for the dihydrates and four steps for m-CPPT β. A third step occurred in the range 300-400 °C for t-CPPD and 275-350 °C for m-CPPD, and this led to a dehydrated phase β-Ca 2 P 2 O 7 . DTA diagrams indicated that these losses corresponded to two endothermic peaks at 315 and 330 °C for t-CPPD and 300 and 305 °C for m-CPPD.

The thermogravimetric analysis of m-CPPT β has been described previously, and there is an initial weight loss of one H 2 O molecule (0.90 H 2 O) at ca. 30 °C owing to the instability of the structural water in the m-CPPT β phase. [22,27] This loss at low temperature was detected during the preliminary plateau of the thermogravimetric analysis at ca. 30 °C. The weight loss between 80 and 110 °C corresponds to two H 2 O molecules, whereas the one between 300 and 450 °C corresponds to one H 2 O molecule and leads to the anhydrous β-Ca 2 P 2 O 7 . However, the different mechanisms were not described in the previous studies.

These dehydration processes are in good agreement with the structure of the m-CPPT β phase, which has three different coordinations of H 2 O molecules. The slight slope between 110 and 300 °C was decomposed by DTA into two exothermic peaks, which were elucidated by FTIR analysis as an internal hydrolysis described by Equation ( 3) and the formation of anhydrous dicalcium phosphate or monetite [Equation ( 4)], identified by XRD and FTIR spectroscopy.

Ca 2 P 2 O 7 •H 2 O Ǟ 2CaHPO 4 (4)
This reaction was followed by a final loss of a water molecule and reconstitution of the P 2 O 7 4-anions to form to the anhydrous phase β-Ca 2 P 2 O 7 [Equation (5)].

2CaHPO 4 Ǟ H 2 O + Ca 2 P 2 O 7 (5)
The reverse reaction occurred at ca. 400 °C and led to the anhydrous phase β-Ca 2 P 2 O 7 . This mechanism of dehydration was observed only with the m-CPPT β phase.

The amorphous phase exhibited a continuous weight loss from 30 to 500 °C with a total loss of 3.87 H 2 O molecules. This hydration is consistent with the values in the range 3.8-4.2 reported by Slater et al., which varied slightly from sample to sample. [24] Exothermic peaks appeared between 110 and 250 °C for the DTA of the amorphous phase, in the same way as m-CPPT β and at the same temperature at which m-CPPD and t-CPPD experienced their first step of dehydration. According to the phenomenon observed with the m-CPPT β phase, the exothermic peak could correspond to a partial internal hydrolysis at ca. 220 °C, as reported by Slater et al. [24] The total weight losses are reported in Table 5 and correspond to the hydration expected for the crystalline phases.

Figure 6 shows the different crystal or particle morphologies of each pure CPP phase as observed by SEM. The crystals of the dihydrated phases were acicular with a large heterogeneity in size. The t-CPPD crystals seemed larger than the m-CPPD crystals, in agreement with former observations. [12,13] The measured sizes were within the range 1-30 μm for t-CPPD with a maximum width of 5 μm and 1-25 μm for m- CPPD with a maximum width of 1 μm. The bundle morphologies, which correspond to category two spherulites, are supposed to be caused by growth from a nucleus, possibly a single acicular crystal formed during the first period of the synthesis. [47] m-CPPT β has a faceted plate morphology with a diameter of up to 20 μm. Balic ´-Z ˇunic ´et al. [22] suggested that this morphology was related to the layered structure of this compound.

At higher pH, the m-CPPT β crystals formed had a diameter of ca. 5 μm with irregular shape, and the t-CPPD crystals showed an acicular shape, similar to that of m-CPPD, with fine needles of around 20 μm length. The a-CPP sample contained round particles with diameters of ca. 100 nm.

It has been reported that CPP crystals can have a wide range of sizes and could potentially present different morphologies. Christoffersen et al. [17,18,48] have studied the crystal growth and dissolution of t-CPPD and it presents two different morphologies with different dissolution kinetics. This variety of morphologies could partially explain the different inflammatory potentials between the phases and between samples of a phase.

The specific surface area (SSA) of the four hydrated CPP phases synthesised by the method we set up were determined as 1.3 m 2 g -1 for t-CPPD, 5.0 m 2 g -1 for m-CPPD, 4.8 m 2 g -1 for m-CPPT β and 25.7 m 2 g -1 for a-CPP. The SSA values are low for the three crystalline phases, as expected, considering their crystal morphology and size (Figure 6). These values are consistent with data already published. [5,14] In addition to phase structure and morphology, surface area has been suggested as an important factor in the pro-posed mechanism of inflammation based on the rupture of lysosome phospholipid membranes by pyrophosphate species on the surface of the crystals. [3][4][5]

Conclusions

The protocol we established allows a fast one-step synthesis of four phases of CPP of biological interest by monitoring the pH and temperature during the synthesis. t-CPPD and m-CPPD, the two phases detected in the joints of arthritic patients, and m-CPPT β and a-CPP, in vitro precursors of the dihydrates, were obtained without intermediates.

The in vitro study of the formation of synthetic and wellcharacterised hydrated calcium pyrophosphate phases available in large amounts could allow a better understanding of the presence of these phases in vivo in joints and contribute to the clarification of the mechanism by which CPP crystals form in vitro and in vivo. Moreover, fine characterisation of the different synthetic CPP phases could improve their detection in patients suffering from calcium salt crystal diseases.

Experimental Section

General: The hydrated calcium pyrophosphates were synthesised by double decomposition of a potassium pyrophosphate solution and a calcium nitrate solution mixed into a buffer solution at a controlled temperature. The reagents were stored in a dry place without evidence of degradation after 6 months.

Materials

All solutions were prepared with deionised water.

Syntheses: A buffer solution was prepared by adding acetic acid (12 mL) and a variable amount of ammonia solution to water (400 mL). The pH was chosen in a range from 3.6 to 5.8. The solution was stirred and heated to a chosen temperature in the range 25-90 °C (Ϯ2 °C).

Calcium and pyrophosphate reagent solutions were prepared separately by dissolving Ca(NO 3 ) 2 •4H 2 O (7.16 g, 3.03 ϫ 10 -2 mol) and K 4 P 2 O 7 (5.00 g, 1.51 ϫ 10 -2 mol) in deionised water (200 mL). The reagent solutions were then added simultaneously into the buffer solution by using a peristaltic pump at a constant volumetric flow rate (4.5 mL min -1 ) to maintain the stoichiometry of the system.

The first white precipitate was observed from 3 to 10 min after the beginning of the addition of the reagent solution, depending on the synthesis conditions. The complete addition of the reagent solu-tions lasted 45 min. After the addition of the reagent solutions was completed, the precipitate formed was left in the mother solution for 10 min at the precipitation temperature.

Filtration was performed by using a Büchner funnel. The precipitate was then washed three times with deionised water and finally dried in an oven at 37 °C overnight. The powder samples were kept at room temperature in sealed vials. No evidence of phase alteration was observed after more than 1 year of storage under these conditions. The yield for this synthesis method was above 90 % for the crystalline phases (m-CPPD 4.08 g, t-CPPD 4.13 g, m-CPPT β 4.61 g) and above 85 % for a-CPP (4.21 g).

The influence of the concentration of the reagents was studied by varying the amount of potassium pyrophosphate in the reagent solution from 2.5 to 10.0 g in water (200 mL, 37.5 to 150.0 mm, respectively).

Characterisation of Products: Powder X-ray diffraction (XRD) measurements were performed with a Seifert XRD-3000 TT diffractometer with Cu-K α radiation (Cu-K α1 λ = 1.54060 Å and Cu-K α2 λ = 1.54443 Å) and equipped with a graphite monochromator. The XRD patterns were obtained in the 2θ range 2-70°with a step size of 0.02°and a scan step time of 16 s at 298 K.

The crystallographic lattice constants were determined by the Jana2006 software (01/07/2011 version). [START_REF] Petricek | The Crystallographic Computing System[END_REF] A pseudo-Voigt shape function was assumed and the background was determined manually. The refinement of the parameters was performed in the order scale factor, zero shift, cell parameters, profile parameter and asymmetry parameter. The total number of variables refined was nine.

The FTIR experiments were performed with a Thermo Nicolet 5700 FTIR spectrometer. The spectra were recorded in the 4000-400 cm -1 wavelength range with a 64 scan accumulation and 4 cm -1 resolution with powder samples in KBr pellets (2 mg of sample in 300 mg of KBr). The peak positions were determined by using second derivation achieved by the Savitzky-Golay routine with a ninepoint third-order polynomial smoothing. [START_REF] Savitzky | [END_REF] Raman microspectroscopy analysis was performed with a Horiba Jobin Yvon Labram HR800 confocal microspectrometer equipped with an AR-diode laser (λ = 532 nm) over a wavenumber range of 55-1750 cm -1 .

Thermogravimetric analyses (TGA-DTA) were performed with a Setaram Instrumentation Setsys evolution system from 30 to 500 °C at a heating rate of 5 °C per minute.

Scanning electron microscopy (SEM) observations were performed with a Leo 435 VP microscope. The accelerating voltage was fixed at 7 kV. Samples were silver plated before observation.

A standard spectrophotometric determination of the yellow phosphovanadomolybdic acid complex was used to determine phosphate concentration. Absorption measurements were performed with a Hitachi U-1100 spectrometer set at 460 nm. The pyrophosphate concentrations were determined after hydrolysis of the pyrophosphate ions into phosphate ions. The samples were hydrolysed at 100 °C in acidic medium for 1 hour, and then the phosphate ions were titrated as described above.

The calcium concentrations were determined by complexometry with ethylenediaminetetraacetic acid (EDTA). [START_REF] Charlot | Chimie Analytique Quantitative[END_REF] The calcium and phosphate titrations had estimated standard deviations equivalent to 1 % of the reading.

The specific surface areas of the samples were evaluated with a Quantachrome Instruments Monosorb Nova 1000 by using the Brunauer-Emmett-Teller method (nitrogen adsorption) with an estimated standard deviation of 0.5 m 2 g -1 .

  [a] CIRIMAT, INPT-CNRS-UPS, Université de Toulouse, ENSIACET, 31030 Toulouse, France E-mail: christele.combes@ensiacet.fr [b] CIRIMAT, INPT-CNRS-UPS, Université Paul Sabatier, 31062 Toulouse, France

Figure 1 .

 1 Figure 1. Synthesis domains of the four pure hydrated CPP phases. Mixtures of these phases were obtained under conditions outside these domains.

Figure 2 .

 2 Figure 2. XRD patterns of the hydrated calcium pyrophosphate phases synthesised.

FTIR

  and Raman spectroscopy allow clear identification of pyrophosphate compounds and provide valuable information on the confirmation of the pyrophosphate ion. Thus, thanks to the flexibility of the pyrophosphate anion in the structure, these methods provide an effective means to clearly identify calcium pyrophosphate hydrates. Figures 3 and 4 show FTIR and Raman spectra, respectively, of the four as-synthesised hydrated CPP phases; the positions and assignments of the IR and Raman bands for the three crystalline hydrated CPP phases are reported in Table

Figure 3 .

 3 Figure 3. FTIR spectra of the four hydrated calcium pyrophosphate phases synthesised. (a) 4000-400 cm -1 domain, (b) 1300-500 cm -1 domain. Although Cornilsen predicted 21 Raman and 21 IR active modes of P 2 O 74-for t-CPPD,[40] these vibrations could not be separated in the spectra, and the data are discussed with regard to the vibrational domains of the free

Figure 4 .

 4 Figure 4. Raman spectra of the four hydrated calcium pyrophosphate phases synthesised. (a) 400-4000 cm -1 domain, (b) 500-1300 cm -1 domain.

Figure 5 .

 5 Figure 5. (a) TGA and (b) DTA curves of the pure hydrated calcium pyrophosphate phases synthesised. Weight losses corresponding to approximately half a H 2 O molecule (0.45 H 2 O) and one H 2 O molecule (1.11 H 2 O) were observed for t-CPPD and m-CPPD, respectively, between 30 and 275 °C. Two steps were noted in this range, and they were correlated to wide endothermic peaks on the differential thermal analysis diagram for both compounds.

Figure 6 .

 6 Figure 6. SEM micrographs of the four synthesised hydrated calcium pyrophosphate phases. (a) t-CPPD, (b) m-CPPD, (c) m-CPPT β and (d) a-CPP.

Table 1 .

 1 Synthesis temperature and pH of the selected samples further characterised by powder XRD and FTIR spectroscopy ([Ca 2+ ] = 2[P 2 O 74-] = 150 mm).

		a-CPP	m-CPPD β m-CPPD t-CPPD
	Temperature [°C] 25	25	90	90
	pH	5.8	4.5	5.8	3.6

Table 3 .

 3 Vibrational band position [cm -1 ] assignments and calculated P-O bond lengths for m-CPPD, t-CPPD and m-CPPT β.

  and 137.8 (two sites), 134.1 and 123°, respectively. The configuration of the P 2 O 7 4-anion in m-CPPD was determined to be closer to those of anhydrous or tetrahydrated compounds than to that of the other dihydrate, t-CPPD, or even other types of dihydrates such as Mn 2 P 2 O 7 •2H 2 O, Fe 2 P 2 O 7 •2H 2 O, Co 2 P 2 O 7 •2H 2 O and Mg 2 P 2 O 7 •2H

  2

  •2H 2 O, Fe 2 P 2 O 7 •2H 2 O, Co 2 P 2 O 7 •2H 2 O and Mg 2 P 2 O 7 •2H 2 O present a dichromate structure with O-P•••P-O torsion angles and P-O-P angles between those of t-CPPD and α-Ca 2 P 2 O 7 .

Table 4 .

 4

	Phosphorus (as P 2 O 7 4-, PO 4	3-or HPO 4 2-ions) and cal-
	cium contents in the four synthesised hydrated CPP phases deter-
	mined by chemical analyses.
		P as P 2 O 7 4-[wt.-%]	P as PO 4 3-or HPO 4 2-[wt.-%]	Ca as Ca 2+ [wt.-%]	Ca/P
	t-CPPD	21.2	0.29	27.6	1.01
	m-CPPD	21.2	0.57	27.6	1.01
	m-CPPT β	19.1	0.23	24.6	1.00
	a-CPP	19.5	0.86	25.2	1.00

Table 5 .

 5 Total weight loss in the pure hydrated CPP phases synthesised. Including loss of water during the preliminary plateau at 30 °C: a-CPP 0.31 H 2 O and m-CPPT β 0.90 H 2 O.

	t-CPPD	m-CPPD m-CPPT β [a] a-CPP [a]
	H 2 O per unit 2.12	2.14	3.73	3.87
	[a]

:

  Anhydrous tetrapotassium pyrophosphate (K 4 P 2 O 7 ) was prepared by heating dipotassium phosphate (K 2 HPO 4 , VWR, Analytical reagent, 100 g) at 400 °C for 3 h in a muffle furnace.

	Anhydrous tetrapotassium pyrophosphate was analysed by XRD
	and FTIR spectroscopy. No detectable amount of dipotassium
	phosphate was observed, and the FTIR spectrum showed the char-
	acteristic vibrational modes of P 2 O 7 4-ions (data not presented).
	Calcium nitrate tetrahydrate salt [Ca(NO 3 ) 2 •4H 2 O, Carlo Erba,
	ACS], ammonia solution (NH 4 OH, VWR, 30 %) and acetic acid
	(C 2 H 4 O 2 , VWR, NormaPur) were used as received without further
	purification.

Acknowledgments

The authors thank the Institut National Polytechnique de Toulouse (PRECIPYCA project -BQR INPT 2011), the Centre National de la Recherche Scientifique (CNRS) (CalArthros project -"Longévité et Vieillissement 2010" CNRS interdisciplinary program) and the Agence Nationale de la Recherche (ANR) (CAPYROSIS project -ANR-12-BS08-0022-01) for supporting this research work.