Katarína Šútorová

Lubomír Prokeš

Virginie Nazabal

Emeline Baudet

Josef Havel

Petr Němec
email: petr.nemec@upce.cz

Laser desorption ionisation time-of-flight mass spectrometry of glasses and amorphous films from Ge-As-Se system

Laser Desorption Ionization Time-of-Flight Mass Spectrometry was exploited for the characterization of Ge-As-Se chalcogenide glasses and corresponding thin films fabricated using pulsed laser deposition. Main achievement of the paper is the determination of laser generated clusters' stoichiometry. The clusters observed were Asb + (b = 1-3), Se2 -, binary

GeAs3Se2 -, Ge2As2Se2 -, Ge2AsSec -(c = 6-7), and GeAs3Sec -(c = 5-6) (in positive as well as in negative ion mode). The stoichiometries of identified species are compared with the structural units of the glasses/thin films revealed via Raman scattering spectra analysis. Some species are suggested to be fragments of bulk glass as well as thin films. Described method is useful also for the evaluation of the contamination of chalcogenide glasses or their thin films.

I. Introduction

Chalcogenide glasses and their thin films form an important class of amorphous materials with applications covering optical fibers, optical waveguides, sensors, phase-change memories, etc. Particularly interesting for optical applications is the good transparency of chalcogenide glasses in the mid-infrared spectral region combined with their high linear and nonlinear refractive indices. [1][START_REF] Eggleton | Chalcogenide photonics[END_REF][START_REF] Lucas | Telluride glasses for far infrared photonic applications[END_REF] Various techniques are known and used to synthesize the amorphous chalcogenides in different forms like bulk glasses, optical fibers, and thin films. 1,[START_REF] Adam | Chalcogenide Glasses for Infrared Photonics[END_REF] Among the chalcogenide glasses, the Ge-As-Se ternary system presents some attractive characteristics, such as one of the broadest glass-forming regions of any of the chalcogenides, relatively easy fabrication in A c c e p t e d m a n u s c r i p t bulk form, wide range of glass transition temperature (Tg), and high nonlinear refractive index. [START_REF] Borisova | Glassy Semiconductors[END_REF][START_REF] Bulla | On the properties and stability of thermally evaporated Ge-As-Se thin films[END_REF] . Physical properties of bulk Ge-As-Se glasses were heavily reported: linear and nonlinear optical properties, [START_REF] Wang | Systematic z-scan measurements of the third order nonlinearity of chalcogenide glasses[END_REF][START_REF] Zha | Optical properties and structural correlations of GeAsSe chalcogenide glasses[END_REF][START_REF] Prasad | Properties of GexAsySe1-x-y glasses for all-optical signal processing[END_REF][START_REF] Gopinath | Third order nonlinearities in Ge-As-Se-based glasses for telecommunications applications[END_REF][START_REF] Quemard | Chalcogenide glasses with high non linear optical properties for telecommunications[END_REF] structure, [START_REF] Zha | Optical properties and structural correlations of GeAsSe chalcogenide glasses[END_REF][START_REF] Prasad | Properties of GexAsySe1-x-y glasses for all-optical signal processing[END_REF][START_REF] Wang | Investigation of the structure of GexAsySe1-x-y glasses by x-ray photoelectron spectroscopy[END_REF] photoinduced processes, [START_REF] Calvez | Light-induced matrix softening of Ge-As-Se network glasses[END_REF] etc. Fabrication and characterization of amorphous Ge-As-Se thin films was subject of deep research in last decade; pulsed laser deposition (PLD), [START_REF] Jarvis | Thin film deposition of Ge33As12Se55 by pulsed laser deposition and thermal evaporation: Comparison of properties[END_REF][START_REF] Zakery | Nonlinear optical properties of pulsed-laserdeposited GeAsSe films and simulation of a nonlinear directional coupler switch[END_REF][START_REF] Choi | Fabrication of low loss Ge33As12Se55 (AMTIR-1) planar waveguides[END_REF][START_REF] Nemec | Photo-stability of pulsed laser deposited GexAsySe100-x-y amorphous thin films[END_REF][START_REF] Hawlova | Photosensitivity of pulsed laser deposited Ge20As20Se60 and Ge10As30Se60 amorphous thin films[END_REF] thermal evaporation [START_REF] Bulla | On the properties and stability of thermally evaporated Ge-As-Se thin films[END_REF][START_REF] Jarvis | Thin film deposition of Ge33As12Se55 by pulsed laser deposition and thermal evaporation: Comparison of properties[END_REF][START_REF] Su | The dependence of photosensitivity on composition for thin films of GexAsySe1-x-y chalcogenide glasses[END_REF][START_REF] Yan | Photoluminescence in Er-doped Ge-As-Se chalcogenide thin films[END_REF][START_REF] Khan | Role of Ge : As ratio in controlling the lightinduced response of a-GexAs35-xSe65 thin films[END_REF][START_REF] Yang | A photo-stable chalcogenide glass[END_REF] or radio-frequency magnetron sputtering [START_REF] Turnbull | Fabrication of waveguides in sputtered films of GeAsSe glass via photodarkening with above bandgap light[END_REF] were shown to be suitable deposition techniques; each of them possessing its own advantages and disadvantages. PLD technique seems to be promising according to its easy control of the deposition process, possibility to fabricate multilayered structures and often stoichiometric material transfer from the target to the films. [START_REF] Nemec | Pulsed laser deposited amorphous chalcogenide and aluminosilicate thin films and their multilayered structures for photonic applications[END_REF][START_REF] Nazabal | Sputtering and Pulsed Laser Deposition for Near-and Mid-Infrared Applications: A Comparative Study of Ge25Sb10S65 and Ge25Sb10Se65 Amorphous Thin Films[END_REF] An important role for the properties of amorphous chalcogenides is played by the presence of impurities. The measured optical losses in chalcogenide glasses are substantially higher than the intrinsic loses, mainly because of the presence of impurities in the glasses. There are three main sources of impurities in chalcogenide glasses: initial elements used for the synthesis, silica glass ampoule usually utilized for the synthesis of glass-forming melt, environment and the residual vacuum gases during the synthesis and treatment of glass/films. 1 For illustration, laser ablation ionization of germanium which is typically used for the synthesis of Ge-based chalcogenide glasses, in spite of its high purity (99.9999%), leads to the formation of large number of species; specifically, Gem clusters are accompanied by GemHn ones, as observed by mass spectrometry (MS). [START_REF] Houska | Laser desorption time-of-flight mass spectrometry of atomic switch memory Ge2Sb2Te5 bulk materials and its thin films[END_REF] It has been demonstrated that Laser Desorption Ionization Time-of-Flight Mass Spectrometry (LDI TOFMS) is a strong experimental method for the generation and identification of clusters of diverse solid state materials inclusive of amorphous chalcogenides as well as for the assessment of their contamination (oxygen, hydrogen and/or carbon impurities). [START_REF] Pangavhane | Laser desorption ionization time-of-flight mass spectrometric study of binary As-Se glasses[END_REF][START_REF] Pangavhane | Laser desorption ionization time-of-flight mass spectrometry of erbiumdoped Ga-Ge-Sb-S glasses[END_REF] In present work, Ge-As-Se glasses and PLD thin films were analysed through LDI

II. Experimental procedures

Chalcogenide glasses from Ge-As-Se system (with Ge10As30Se60, Ge10As35Se55, Ge10As40Se50, Ge15As30Se55, and Ge20As20Se60 nominal compositions) were synthesized from high purity elements (Ge, As, and Se, 5-6 N) using the conventional melting and quenching technique. [START_REF] Nemec | Photo-stability of pulsed laser deposited GexAsySe100-x-y amorphous thin films[END_REF] For the fabrication of thin films, UV PLD was employed using 25 mm targets made from above described bulk glasses, PLD setup consisting of a KrF excimer laser (248 nm, 300±3 mJ per pulse, 30 ns pulse duration, 20 Hz repetition rate), and a vacuum chamber (background pressure <3x10 -4 Pa). The substrates used for PLD (chemically cleaned microscope glass slides, Si wafers) were positioned parallel to the target surface at the target to substrate distance of 5 cm. The energy fluence on the target was set at ~2.6 J.cm -2 . Both the target and the substrates were rotated in off-axis geometry in order to avoid deep damage of the targets and to improve the thickness homogeneity of the films.

A diffractometer with Bragg-Brentano θ-θ geometry (D8-Advance, Bruker AXS, Karlsruhe, Germany) was employed for the recording of X-ray diffraction (XRD) data. The diffraction angles were measured at room temperature from 5 to 65° (2θ) in 0.02° steps with a counting time of 5 s per step. To determine chemical composition of fabricated Ge-As-Se glasses/films, a scanning electron microscope equipped with an energy-dispersive X-ray analyzer (EDX, JSM 6400-OXFORD Link INCA, JEOL Ltd., Tokyo, Japan) was used. Ge-As-Se materials were studied also using MIRA3 SEM (TESCAN, Brno, Czech Republic).

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The Raman scattering spectra of Ge-As-Se samples were measured at room temperature by LabRAM HR spectrometer (Horiba Jobin-Yvon, Villeneuve d'Ascq, France) with 785 nm laser as excitation source. The intensity of laser beam was kept at low level to avoid thermally induced structural transformations induced by absorption of high laser power densities. The frequency resolution of used configuration was ~0.7 cm -1 per pixel.

Optical functions (refractive index and extinction coefficient spectral dependencies) of bulk glasses/thin Ge-As-Se films (and thicknesses in case of thin films) were obtained from the analysis of data measured using an ellipsometer with automatic rotating analyzer (VASE, J. A. Woollam Co., Inc., Lincoln, NE, USA). The measurement parameters were following: spectral region 300-2300 nm with wavelength steps of 10 nm (20 nm in case of target glasses), angles of incidence 50°, 60°, and 70°. For the analysis of VASE data, Cody-Lorentz (CL) model was used. [START_REF] Cody | Semiconductors and semimetals[END_REF] LDI TOFMS data were recorded within the powdered bulk glasses as well as corresponding thin films. The powders of Ge-As-Se glasses were prepared exploiting vibratory ball mill with agate mortar and grinding ball. For the mass spectrometry, 1 µL of acetonitrile suspension of the powder (1 mg/mL) was deposited on the target plate and dried in a stream of air at room temperature. The target plate was always cleaned carefully with ethanol and double distilled water before the samples deposition. [START_REF] Sutorova | Laser Desorption Ionisation Time-of-Flight Mass Spectrometry of Chalcogenide Glasses from (GeSe2)100-x(Sb2Se3)x System[END_REF] AXIMA Resonance mass spectrometer equipped with a quadrupole ion trap mass analyser (Kratos Analytical Ltd., Manchester, UK) employing a time-of-flight mass analyser was used for the measurement of LDI TOFMS data. Using AXIMA Resonance device, MS signal was recorded in m/z ranges of 100-400, 250-1200, 800-3500, 1500-8000, and 3000-15000. The instrument used a delayed extraction. All the measurements were performed in both positive and negative reflectron ion modes. For necessary external mass calibration, a calibration standard (red phosphorus) was used in both ionisation modes. [START_REF] Sutorova | Laser Desorption Ionisation Time-of-Flight Mass Spectrometry of Chalcogenide Glasses from (GeSe2)100-x(Sb2Se3)x System[END_REF] For LDI TOFMS, a UV

A c c e p t e d m a n u s c r i p t source (N2 laser operating at 337 nm) with 3 ns pulse duration and repetition rate of 5 Hz was applied. The laser energy was scaled in arbitrary units (a.u.) from 0 to 180; this relative scale will be used hereinafter. Laser spot size on sample had circular shape with ~150 μm diameter.

Resulting laser energy fluence was ~1 J.cm -2 . At least 1500 laser pulses were applied for the accumulation of each mass spectrum. Experimental isotopic patterns were compared with theoretical ones which were calculated through Launchpad software (Kompact version 2.9.3, 2011, Kratos Analytical Ltd., Manchester, UK).

III. Results and discussion

Fabricated glasses from Ge-As-Se system and corresponding thin films prepared by PLD were found to be amorphous and homogeneous; the films are without cracks/defects, according to optical/electron microscopy and XRD patterns. The morphology of the layers is of good quality as results from SEM micrographs. The chemical composition of Ge-As-Se glasses/thin films, as determined using EDX when averaging several measured points, is in good agreement with nominal one. In fact, all the measured compositions (for both bulks and films) are within ~3 at. % variation of the nominal composition based on quantities of the elements used to prepare the glasses. Taking into account error limits of the EDX technique (±1 at. %), one can conclude that the material transfer during PLD is close to target stoichiometry.

The thicknesses of all the films varied from ~960 to ~1230 nm, as determined by VASE data analysis. Obtained refractive indices spectral dependences for both, the bulk glasses and PLD thin films are illustrated in Fig. 1.; the data presented here show relatively good agreement between refractive indices of bulk glasses and thin films, especially in nearinfrared spectral region. The CL model reliability for the analysis of VASE data is confirmed by low values of mean square error of the fitting procedure (<7.2).

A c c e p t e d m a n u s c r i p t

Mass spectra of Ge-As-Se samples were measured aiming at the stoichiometry determination of singly positively and negatively charged GeaAsbSec ions formed through LDI. Stoichiometry of species contained in the plasma was evaluated based on the comparison of experimental isotopic patterns with the theoretical ones. We note that the atomic weights of Ge, As, and Se are close to each other (72.64, 74.92, and 78.94 g.mol -1 for Ge, As, and Se), therefore the isotopic envelopes are complex when number of Ge, As or Se atoms are combined.

Firstly, we deal with the LDI TOFMS of bulk Ge-As-Se glasses. Here, we used similar methodology as the one we successfully applied earlier. [START_REF] Sutorova | Laser Desorption Ionisation Time-of-Flight Mass Spectrometry of Chalcogenide Glasses from (GeSe2)100-x(Sb2Se3)x System[END_REF] The influence of laser energy on TOFMS spectra was studied in negative ion mode. We have found that the ionization started when the laser energy reached ~80 a.u. At higher laser energies, intensities of the LDI TOFMS signal gradually climbed. This observation is presented in Fig. 2. Sufficient mass resolution of the mass spectra and highest number of detected clusters in negative ion mode were identified at the laser energy ~120-130 a.u.

The MS spectrapeak with lowest mass was assigned to Se2 -ion. The further observed peaks were recognized as AsSec -(c = 2-3) and As3Se2 -clusters. The GeAsSe4 -cluster was found to overlap with binary As3Se3 -one (71 and 29 %). Percentage in brackets means the contribution of individual species to the overall peak intensity. Few oxidized clusters were also determined as SeAsO -, SeAsO2 -, Se2AsO -, GeAsSeO11 -, and Ge3AsSeO3 -. However, mentioned oxidized clusters were observed with low intensity and their content can be estimated below < 0.001 % rel. Unfortunately, precise quantitative evaluation of oxygen containing species content is impossible as the exact correlation between the measured quantity (MS signal) and the number of generated ions is not known.

In addition to above mentioned clusters, also several lower intensity clusters were found.

The overlap of binary Ge2As2 -and Ge3As -clusters (83 and 17 %) was observed. In higher We turn to the analysis of Ge-As-Se thin films. In case of LDI TOFMS analysis of Ge-As-Se PLD thin films performed in negative ion mode, the ionization began when the energy of the laser reached ~100 a.u. Also here, at higher laser energies, intensities of the observed clusters gradually increased. At the same laser energy as for bulk glasses (~120 a.u), mass spectra containing highest number of ionized species in negative ion mode were recorded.

The lowest mass peak was identified as Se2 -ion. The further observed peaks were recognized as AsSec -(c = 2-3) and As3Se2 -clusters. The GeAsSe4 -cluster was found to overlap with binary As3Se3 -one (71 and 29 %). In addition to above mentioned clusters several low intensity clusters were detected. The overlap of Ge2Se2 -with As2Se2 -cluster was found (52 and 48 %). The other observed peaks were identified as overlaps of binary Ge3As2 - one with As2Se3 -and Ge2Se3 -clusters (6, 63 and 31 %); ternary Ge2As2Se2 -one with

A c c e p t e d m a n u s c r i p t

GeAs2Se3 -and GeAs3Se2 -clusters (9, 27 and 64 %), and binary Ge3Se4 -cluster with As3Se4 - one (33 and 77 %). No more clusters were found for m/z values over 600. It is worth noting that in contrast with the MS analysis of bulk glasses, no clusters containing oxygen were detected in Ge-As-Se thin films. This difference could be explained by the procedure of MS samples preparation. The films were processed by PLD under vacuum. In case of bulk samples, the powders were made using vibratory ball mill without protecting inert atmosphere, which may lead to the slight surface oxidation. As we expect that oxidation proceeds in very top layer of the chalcogenide powder particles only, the presence of oxidized species should be marginal. This is confirmed by MS spectra; estimated content of oxidized clusters is below < 0.001 % rel. as mentioned above.

When analyzing Ge-As-Se thin films in positive ion mode, impact of varying laser energy was also investigated. Similarly to negative ion mode, the ionization started at the laser energy ~100 a.u.. At higher laser energies, intensities of the observed clusters gradually increased (Fig. 4). Mass spectra in positive ion mode present adequate resolution, signal intensity, and highest number of detected species after reaching ~120-130 a.u. laser energy.

The MS spectra peaks were identified as AsbSe + (b = 1-3) clusters and As3Se2 + one as for the analysis of bulk Ge-As-Se glasses. Also, arsenic clusters Asb + (b = 1-3) were identified.

For reader's convenience, a summary of species identified in mass spectra of plasma generated from all Ge-As-Se samples is given in Table I. It is worthy to point out that some species summarized in Table I. are common for both types of samples (bulk glasses and thin films); these are AsbSe + (b = 1-3), As3Se2 +/-, Se2 -, AsSec -(c = 2-3), GeAsSe4 -, and As3Se3 -.

On the basis of Raman scattering spectra (Fig. 5), the structure of Ge-As-Se bulk glasses can be determined. The local structure is formed by different structural motifs as described below. First, the presence of GeSe4 corner-sharing tetrahedra is documented through Raman bands peaking at ~193-198 and ~300 cm -1 originating from A1(ν1) symmetric stretching and Raman band peaking at ~215 cm -1 is attributed to the "companion" mode due to the vibrations of GeSe4 tetrahedra sharing their edges. Homopolar Ge-Ge bonds can also be presented [START_REF] Nemec | Photo-stability of pulsed laser deposited GexAsySe100-x-y amorphous thin films[END_REF] in Ge-GemSe4-m, m=1,2,3,4 structural units with wide Raman band centered at ~270 cm -1 and a shoulder at ~170 cm -1 . Arsenic based structural motifs are represented by AsSe3 pyramids identified via stretching vibration modes at ~224-229 cm -1 . The presence of cage-like As4Se3(4) molecules containing homopolar As-As bonds can be connected with the Raman features at ~238, 196, 256, 248, and 190 cm -1 . [START_REF] Nemec | Photo-stability of pulsed laser deposited GexAsySe100-x-y amorphous thin films[END_REF][START_REF] Yang | Correlation between structure and physical properties of chalcogenide glasses in the AsxSe1-x system[END_REF] The Raman spectra of as-deposited Ge-As-Se layers show similar features as the bulk glasses do (Fig. 5). The widths of individual bands are generally larger for the thin films when compared with bulk glasses, which is an indication of a larger degree of disorder in the films.

The local structure of Ge-As-Se films is, in accordance with parent bulk glasses, formed by GeSe4 corner-and edge-sharing tetrahedral and arsenic-based structural units as discussed below.

The differences between Raman spectra of bulk Ge-As-Se glasses and corresponding thin films concern mainly ~210-260 cm -1 region for glasses/films containing lower amount of selenium (Ge10As35Se55, Ge10As40Se50, Ge15As30Se55). Specifically, in case of Ge-As-Se films with lower selenium content, the Raman bands in ~210-260 cm -1 region are almost featureless. This behaviour was already reported for As-Se PLD films, where for selenium under-stoichiometric compositions (As50Se50, As57Se43, As60Se40), amorphous layers present featureless character of Raman response in discussed spectral range while for bulk glasses several bands can be distinguished. [START_REF] Nemec | Structure, optical properties and their photo-induced changes in AsxSe100-x (x=50, 57.1, 60) amorphous thin films prepared by pulsed laser deposition[END_REF] Further, for As40Se60 stoichiometric composition, Raman spectra of the bulk glass and as-deposited PLD films are almost identical, again featureless. [START_REF] Nemec | Structure, Thermally and Optically Induced Effects in Amorphous As2Se3 Films Prepared by Pulsed Laser Deposition[END_REF] Moreover, the shape of Raman spectra of as-deposited PLD As-rich As-Se layers is similar to stoichiometric As40Se60 glasses/PLD films. Observed Raman spectra of As-rich As-Se bulk glasses/PLD films are interpreted in terms of local structure differences between the glasses and films. The structure of As-rich As-Se bulk glasses is formed by AsSe3 pyramids, cagelike molecular units (As4Se3(4)), and arsenic clusters. In contrast, the structure of corresponding as-prepared PLD thin films is probably formed mainly by As-based pyramids (AsSe3, As2Se2, As3Se, and eventually As4) with vibrational frequencies near the frequencies of AsSe3 pyramids. [START_REF] Nemec | Structure, optical properties and their photo-induced changes in AsxSe100-x (x=50, 57.1, 60) amorphous thin films prepared by pulsed laser deposition[END_REF] Nevertheless, presence of some As4Se3(4) cage-like molecules in the structure of Ge-As-Se PLD films is also probable. We believe that analogous interpretation is valid also for studied Ge-As-Se glasses/PLD films with lower selenium content.

To conclude, glassy network of Ge-As-Se glasses and amorphous layers is mainly formed by GeSe4 tetrahedra and different arsenic-based structural units. In accordance with the chemical composition of the samples, homopolar Ge-Ge and As-As bonds are also expected.

The results of LDI TOFMS analysis may bring proof for some of above mentioned structural units. For instance, AsSe3 -entities obviously validate the presence of AsSe3 pyramidal units in the structure of Ge-As-Se glasses and/or thin films. Evidence of AsbSe + (b = 1-3) and As2Se2 - species is coherent with the presence of arsenic-based pyramids (AsSe3, As2Se2, As3Se).

As3Se2 +/-, As3Se3 -, and As2Se3 -clusters are probably connected with the fragmentation of As4Se3(4) cage-like molecules. The fragmentation of As4Se3(4) cage-like molecules and/or arsenic-based pyramids can also lead to the presence of Se2 -and AsSe2 -species in the plasma.

Identification of Asb + (b = 1-3) species can be attributed to the fragmentation of arsenic-rich As-Se pyramids, whose presence is expected in the structure of Ge-As-Se films.

Contrary, in case of GeSe4 tetrahedra, which constitute one of the most significant structural entities of Ge-As-Se glasses/films as proved by Raman data, LDI TOFMS did not detect clear evidence of such structural units. The LDI TOFMS data recorded in this work

show much lower number of germanium containing species than would be anticipated taking into account the stoichiometry of materials under study. This observation might be connected with difficult ionization of Ge-based species. [START_REF] Pangavhane | Laser desorption ionization time-of-flight mass spectrometry of erbiumdoped Ga-Ge-Sb-S glasses[END_REF] Some binary (such as Ge2As2 -, Ge3As -, Ge2As3 -, Ge3As2 -, As5Se4 -, As5Se5 -, Ge2Se2 -, Ge2Se3 -, Ge3Se4 -or As3Se4 -one) and ternary species identified in mass spectra of Ge-As-Se glasses and thin films (Table I.) do not agree with structural motifs found by the analysis of Raman spectra of Ge-As-Se glasses/thin films. The presence of unexpected species identified by LDI TOFMS likely originates from deeper fragmentation of basic structural units. Due to very high energy of species forming plasma plume, one should take into account also reactions ongoing in plasma state.

Unfortunately, quantitative correlation between Raman data and LDI TOFMS results is difficult to obtain. This is mainly due to the fact that precise correlation between measured mass spectrometry signal and the number of generated ion species is not known. Further problem is the fragmentation of the glass structure which is probably complex process and the same fragments can be formed from different glass structural units or different fragments can be formed from the same structural unit. The fragmentation can also be different for varying glass composition.

To conclude, the species which originate from the interaction of Ge-As-Se glasses with high energy laser pulses partially correspond with the structural units observed in original material. This conclusion can be supported by the fact that the stoichiometry of the species identified through LDI TOFMS does not vary dramatically with increasing laser energy.

IV. Conclusions

In this paper, refractive index (through spectroscopic ellipsometry) and structure (via Raman scattering spectroscopy) of Ge-As-Se bulk glasses and pulsed laser deposited thin films were studied. Refractive indices of bulk glasses and thin films are in satisfactory agreement in nearinfrared spectral region. In case of low Se content, the structure of the glasses is partly

, c = 1-4), Ge2Sec -(c = 2-3), Ge2Asb -(b = 2-3), Ge3Asb -(b = 1-2), Ge3Se4 -, As5Sec -(c = 4-5), GeAsSe4 -, GeaAsSe5 -(a = 1-4), GeAs2Se3 -, GeAs3Se2 -, Ge2As2Se2 -,
Ge2AsSec -(c = 6-7), and GeAs3Sec -(c = 5-6) entities were identified. Some identified species were observed for both, bulk glasses and thin films. LDI TOF MS showed also the presence of some oxygen-containing impurities in the studied bulk materials. Some of clusters such as AsSe3 -, AsbSe + (b = 1-3), and As2Se2 -detected in mass spectra suggest to be species present in the structure of original glasses as well as thin films fabricated by pulsed laser deposition.

Other charged clusters detected in the plasma plume do not seem to be involved in the structure of Ge-As-Se glasses/films; such species are probably created through the reactions of highly energetic species in the plasma plume or due to deeper fragmentation of basic structural units. Laser desorption ionization time-of-flight mass spectrometry technique confirmed high purity of synthesized Ge-As-Se glasses and thin films as only few species containing oxygen were identified.

A c c e p t e d m a n u

  s c r i p t TOFMS in order to acquire information regarding the structural fragments (and their contamination) present in the plasma emerging from the interaction of pulsed laser with solid state material. The obtained results would contribute to a deeper comprehension of the preparation and properties of amorphous chalcogenide thin films fabricated by pulsed laser deposition and generally via plasma deposition techniques.

A c c e p t e d m a n u s c r i p t

  mass range (above 300 m/z), other entities take place. These were established as overlap of Ge2As3 -and Ge3As2 -species (46 and 57 %). In addition, ternary GeaAsSe5 -(a = 1-2) clusters were identified. The heaviest clusters were found as overlap of ternary clusters with binary one for m/z values ~690 and ~772; Ge2AsSe6 -one with GeAs3Se5 -and As5Se4 -clusters (43, 22 and 35 %), Ge2AsSe7 -one with GeAs3Se6 -and As5Se5 -clusters (45, 18 and 37 %). Finally, Ge3AsSe5 -and Ge4AsSe5 -species were detected. No more clusters were identified at m/z values over 800. Fig.3presents a comparison between experimental and model mass spectra in m/z range 150-320; we note a perfect agreement.The influence of laser energy on TOFMS spectra was studied also in positive ion mode; the ionization started when the laser energy reached ~80-90 a.u. At higher laser energies, intensities of the observed clusters decreased with rising the energy of the laser. At the laser energy ~120 a.u., mass spectra have satisfactory mass resolution, intensity of signal, and highest number of clusters in positive ion mode. The peaks detected were identified as signal coming from AsbSe + (b = 1-3) clusters and As3Se2 + one.

A c c e p t e d m a n u s c r i p t F2

 t asymmetric vibration mode of GeSe4 corner-sharing tetrahedral, respectively. Second,

A c c e p t e d m a n u s c r i p t

A c c e p t e d m a n u s c r i p t

A c c e p t e d m a n u s c r i p t

  different from PLD films. Apart from GeSe4 tetrahedra, bulk glasses contain AsSe3 pyramids, cage-like molecular units (As4Se3(4)), and arsenic clusters. In contrast, the structure of asprepared PLD thin films is probably formed by As-based pyramids (AsSe3, As2Se2, As3Se, and eventually As4). Laser desorption ionization time-of-flight mass spectrometry was used for the analysis of plasma formed from Ge-As-Se chalcogenide materials. In positive ion mode, Asb + (b = 1-3), AsbSe + (b = 1-3), and As3Se2 + species were found. In negative ion mode, Se2 -, AsbSec -(b = 1-3

A c c e p t e d m a n u s c r i p t Figure captions Fig. 1 .

 captions1 Figure captions

A c c e p t e d m a n u s c r i p tFig. 2 .

 2 Fig. 2. The effect of laser energy on the LDI TOF mass spectra of the bulk Ge10As40Se50 sample showing the formation of As3Se3 -cluster overlapped with GeAsSe4 -one, GeaAsSe5 -(a = 1-2) clusters, and overlap of As5Se4 -one with Ge2AsSe6 -and GeAs3Se5 -clusters.

A c c e p t e d m a n u s c r i p t

Fig. 3 .

 3 Fig. 3. Selected part of the mass spectrum concerning bulk Ge10As30Se60 powdered sample showing the formation of Se2 -ion and AsSec -(c = 2-3) clusters in comparison with theoretical mass spectra. AsSe3 -was magnified 10 times. Conditions: negative reflectron ion mode; laser

A c c e p t e d m a n u s c r i p t

Fig. 4 .

 4 Fig. 4. The effect of laser energy on the LDI TOF mass spectra of the Ge20As20Se60 thin film showing the formation of AsbSe + (b = 1-3) clusters. Conditions: positive reflectron ion mode;

A c c e p t e d m a n u s c r i p tFig. 5 .

 5 Fig. 5. Raman scattering spectra of Ge-As-Se bulk glasses (full curves) and corresponding

*

  Summary of positively or negatively charged species identified in mass spectra of Ge-As-Se bulk glasses and PLD thin films. * For clarity, GeaAsbSec formula convention is used for identification of individual species; c stands for the number of selenium atoms in the clusters. Note that overlapping clusters (isobars) are given inside the { } brackets; in case of overlap of clusters with different c, cluster(s) related to proper c are in boldface.

A c c e p t e d m a n u s c r i p t 5 GeAsSe5

 5

Acknowledgements

Support from the Czech Science Foundation (Project No. 13-05082S) is greatly


Vincent Auffret
email: vincent.auffret@chu-rennes.fr

MD Rishi Puri

PhD Josep Rodés-Cabau

MD Hervé Le Breton

The challenging realm of neurocognitive evaluation following transcatheter aortic valve implantation

Disclosures:

Dr. Auffret received fellowship support from the Fédération Française de Cardiologie, and research grants from Abbott, Edwards Lifesciences, Medtronic, Biosensors, Terumo, and Boston Scientific. Dr. Rodés-Cabau has received research grants from Edwards Lifesciences and Medtronic. Pr Hervé Le Breton has received speaker fees from Edwards Lifesciences and Medtronic. Dr Puri has nothing to disclose.

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Following the first transcatheter aortic valve implantation (TAVI) performed by Cribier and colleagues in 2002 [START_REF] Cribier | Percutaneous Transcatheter Implantation of an Aortic Valve Prosthesis for Calcific Aortic Stenosis: First Human Case Description[END_REF], this procedure has evolved considerably to the point where it is now on the verge of posing a viable treatment option amongst intermediate surgical-risk patients with symptomatic severe aortic stenosis [START_REF] Leon | Transcatheter or Surgical Aortic-Valve Replacement in Intermediate-Risk Patients[END_REF]. In the early days of TAVIs' evolution, cerebrovascular events were considered as the iceberg standing in the way of the "TAVI Titanic" [START_REF] Auffret | Predictors of early Cerebrovascular Events in Patients with Severe Aortic Stenosis Undergoing Transcatheter Aortic Valve Replacement: A Systematic Review and Meta-Analysis[END_REF]. Recent randomized trials reflecting improved patients' selection, increased operators' experience, and iterated devices have somewhat dissipated those fears, with a similar incidence of cerebrovascular events already demonstrated between TAVI and surgical aortic valve replacement [START_REF] Leon | Transcatheter or Surgical Aortic-Valve Replacement in Intermediate-Risk Patients[END_REF][START_REF] Adams | Transcatheter aorticvalve replacement with a self-expanding prosthesis[END_REF]. However, these clinically overt events seem to be the tip of the iceberg, with silent cerebral micro-embolization and its potential cognitive consequences, lurking beneath the surface [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF]. The emergence of embolic protection devices (EPDs), along with the trend towards treating lower surgical-risk patients, shed further light on these specific issues, which are increasingly recognized as a surrogate burden of the ischemic cerebral insult imparted during TAVI [START_REF] Lansky | A prospective randomized evaluation of the TriGuard™ HDH embolic DEFLECTion device during transcatheter aortic valve implantation: Results from the DEFLECT III trial[END_REF][START_REF] Haussig | Effect of a Cerebral Protection Device on Brain Lesions Following Transcatheter Aortic Valve Implantation in Patients With Severe Aortic Stenosis: The CLEAN-TAVI Randomized Clinical Trial[END_REF].

In the setting of surgical aortic valve replacement, a recent review reported rates of 50 to 70% of cognitive decline (CD) within a week of cardiac surgery whereas 10 to 20% of patients exhibited persistent CD at 1 year [START_REF] Patel | Risk Factors Associated with Cognitive Decline after Cardiac Surgery: A Systematic Review[END_REF]. A few seminal studies attempted to navigate the deep and dark waters of cognitive evaluation early post-TAVI, mainly demonstrating preserved, and even improved, cognition [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF]. However, most of these studies focused on global cognition or memory during a short-term follow-up period without serial evaluations. Moreover, changes in cognitive function were assessed by performing formal statistical tests on the mean or median cognitive scores in the setting of small sample-size, therefore implying the risk of type II error. By contrast, recent studies evaluating the use of EPD suggested that rates of early CD, on the basis of the Montréal Cognitive Assessment (MoCA) measurement, may be as high as 50 to 72% at day 2 post-TAVI and 37 to 55% at discharge [START_REF] Lansky | A prospective randomized evaluation of the TriGuard™ HDH embolic DEFLECTion device during transcatheter aortic valve implantation: Results from the DEFLECT III trial[END_REF][START_REF] Haussig | Effect of a Cerebral Protection Device on Brain Lesions Following Transcatheter Aortic Valve Implantation in Patients With Severe Aortic Stenosis: The CLEAN-TAVI Randomized Clinical Trial[END_REF].

Interestingly, the use of EPDs seemed to mitigate this deleterious effect on cognition. Nonetheless, these studies defined CD as a drop of at least 1 point in the MoCA score, thus failing to properly take into account the variability of changes across tests, which depends on the stability and reliability of the assessment tool [START_REF] Wise | Methods for analyzing psychotherapy outcomes: a review of clinical significance, reliable change, and recommendations for future directions[END_REF][START_REF] Frerichs | A comparison of methods for measuring cognitive change in older adults[END_REF]. In other words, in the worst case scenario, these findings only reflect random variability.

To avoid this pitfall of cognitive evaluation, some studies used specific methods to account for this inherent variability [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF][START_REF] Ghanem | Cognitive trajectory after transcatheter aortic valve implantation[END_REF][START_REF] Schoenenberger | Evolution of Cognitive Function After Transcatheter Aortic Valve Implantation[END_REF]. In the first of its type, Ghanem et al. [START_REF] Ghanem | Cognitive trajectory after transcatheter aortic valve implantation[END_REF] evaluated the cognitive trajectory of 111 TAVI recipients, 32 of whom were followed for up to 2 years, using the repeatable battery for the assessment of neuropsychological status (RBANS) with alternate forms to counterbalance practice effects. Post-TAVI CD was defined as a drop of > 1 standard deviation compared with a subject's score pre-TAVI, which occurred in 9% of patients overall. Early CD (within 3 days of TAVI) was apparent in 5.4% of patients, persisting in 50% of them, whereas 3.6% of TAVI recipients suffered from late CD (≥3 months post-TAVI). Only age associated with the occurrence of CD. Among 229 patients ≥ 70 years old undergoing TAVI, Schoenenberger et al. [START_REF] Schoenenberger | Evolution of Cognitive Function After Transcatheter Aortic Valve Implantation[END_REF], using the Mini-Mental State Examination (MMSE) and also defining CD with the use of the baseline standard deviation, reached roughly similar conclusions. In their cohort, 29 patients (12.7%) demonstrated CD (≥3 points decrease in MMSE) at 6 months post-TAVI including 8 patients exhibiting a major decrease (≥5 points) for whom review of their records identified an obvious medical cause. Of note, the authors failed to isolate multivariable predictors of CD. These 2 studies used tests allowing an integrative measure of global cognition through measurements of several cognitive domains. However, the standard deviation method used in these studies may be prone to overclassifying CD compared with more stringent methods such as practice-corrected reliable change index (RCI) or regression based methods [START_REF] Frerichs | A comparison of methods for measuring cognitive change in older adults[END_REF]. Besides, both the MMSE and RBANS lack sensitivity in the detection of mild cognitive impairment. By contrast, the MoCA, albeit less commonly performed than the MMSE, demonstrated a greater sensitivity in the detection of subtle cognitive changes, due to the inclusion of a more comprehensive evaluation of executive functions which are predominantly impaired in vascular cognitive impairment [START_REF] Nasreddine | The Montreal Cognitive Assessment, MoCA: a brief screening tool for mild cognitive impairment[END_REF].

Recently, we reported the cognitive trajectory of 51 TAVR recipients during a 1 year followup period, using the MoCA and practice-corrected RCI to define cognitive changes [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF]. On the basis of the RCI of MoCA score, 4 patients (7.8%) presented with 30-day CD, which persisted at 1 year in 1 patient (2.0%). Overall, 11.8% of patients exhibited CD at 1 year post-TAVI. Using 5 specific tests for some complex executive functions, we demonstrated that a quarter of TAVI recipients experienced a deterioration in at least one of these tests at 1 month post-TAVI, which was transient in 60% of these patients and sustained in 40 % of them (10% of the global cohort). Interestingly, we failed to demonstrate a meaningful association between cognitive changes and subsequent quality-of-life or functional status. Overall, the available evidences suggest that mid-term CD affects ≈10 to 15% of TAVI recipients. Early CD (i.e. the most likely to result from the procedure itself) occurs in 2 to 10% of patients when considering global cognition, may affect 25% of them when specific cognitive domains are evaluated, and seems to persist in 25 to 50% of these patients.

On the contrary, an early and sustained cognitive improvement post-TAVI has also been demonstrated in a sizeable proportion of TAVI recipients [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF][START_REF] Lansky | A prospective randomized evaluation of the TriGuard™ HDH embolic DEFLECTion device during transcatheter aortic valve implantation: Results from the DEFLECT III trial[END_REF][START_REF] Schoenenberger | Evolution of Cognitive Function After Transcatheter Aortic Valve Implantation[END_REF], ranging from 8 to 38% depending on the assessment tool, timing of the evaluation and methods used to define cognitive changes. This improvement seems more likely amongst patients cognitively-impaired at baseline [START_REF] Auffret | Serial Changes in Cognitive Function Following Transcatheter Aortic Valve Replacement[END_REF][START_REF] Schoenenberger | Evolution of Cognitive Function After Transcatheter Aortic Valve Implantation[END_REF]. Importantly, Schoenenberger et al. [START_REF] Schoenenberger | Evolution of Cognitive Function After Transcatheter Aortic Valve Implantation[END_REF] demonstrated a lower pre-TAVI aortic valve area amongst patients with post-procedure cognitive improvement; supporting the hypothesis that improvements in cardiac output and, consequently, in cerebral blood flow post-TAVR may reverse some of the baseline alterations.

Admittedly, the current literature leaves us with as many questions as answers, highlighting the tremendous complexity of cognitive evaluation using heterogeneous methods and definitions. That being said, what should we expect from our scientific journey on the still largely unexplored and winding path of post-TAVI cognition? First and foremost, we urgently need to validate and harmonize a suitable battery of neurocognitive tests for TAVI candidates, particularly within the setting of treating lower surgical-risk patients, especially given the growing interest surrounding neuroprotective strategies such as EPDs. In keeping with the prior point, stringent methods that take into account cognitive tests' variability should be mandatory to define significant cognitive changes post-TAVI in future studies. To fairly evaluate TAVI's role in cognitive evolution, we also need to unravel the specific cognitive trajectory of the elderly population with medically-managed severe aortic stenosis, as this has not been specifically assessed so far. Future research should aim at precisely elucidating the underlying mechanisms of post-TAVI CD. Indeed, most studies so far have focused on the cerebral embolic insult as the leading cause of CD. Although silent cerebral infarcts could conceivably be a major cause of early CD, they may be transient and thus their implication in sustained CD remains questionable. Moreover, only one study demonstrated a relationship between cerebral embolism and the occurrence of CD following surgical aortic valve replacement [START_REF] Barber | Cerebral ischemic lesions on diffusion-weighted imaging are associated with neurocognitive decline after cardiac surgery[END_REF]; a finding that has been recently reproduced in only one small study in the setting of TAVI, demonstrating a moderate but significant correlation between CD and the number and volume of new cerebral lesions on diffusion-weighted magnetic resonance imaging [START_REF] Fanning | Neurological Injury in Intermediate-Risk Transcatheter Aortic Valve Implantation[END_REF]. Identifying predictors of cognitive changes and evaluating neuroprotective strategies in high-risk subgroups are other areas of major interest. Finally, of particular importance is the impact of cognitive changes on TAVI recipients' quality-of-life and functional status; largely unknown to date. There is much still to be discovered in the fascinating land of post-TAVI cognition, and it is safe to say that this incredible journey is thus far from drawing to an end…


K Theuwissen

J Lacaze

L Laffont
email: lydia.laffont@ensiacet.fr

Structure of graphite precipitates in cast iron

This study presents microstructural investigations of graphite precipitates found in cast irons. Binary Fe eC, FeeCeSb and FeeCeCe alloys containing different graphite morphologies (flake and spheroidal) were produced and compared to commercial irons in an attempt to rationalize the effect of Sb and Ce on graphite growth. The extensive use of transmission electron microscopy (mainly electron diffraction and high resolution lattice fringe imaging) enabled further understanding of graphite growth mechanisms. It was found that the inner structure of graphite precipitates consists of growth blocks stacked upon each other, for all investigated morphologies. This suggests that graphite crystals develop mainly by a 2D nucleation and growth mechanism, and that the final shape of the precipitates is associated to the occurrence of crystallographic defects in the graphite lattice (such as twins, misorientations and rotations) during growth.

Introduction

Cast irons are the most important casting materials and their use is justified by the wide range of mechanical properties which can be achieved, associated with their competitive price. These materials have been used for thousands of years, dating back to 500 BC. in ancient China [START_REF] Stefanescu | Solidification and modeling of cast iron-A short history of the defining moments[END_REF]. Cast irons containing iron carbide (cementite) are named "white" whereas those in which carbon is present as graphite are called "grey" irons. In these latter, those having graphite precipitates with a flake-like morphology show a good damping capacity and compressive strength but also low tensile strength and ductility [START_REF]Properties and Selection: Irons, Steels and High Performance Alloys[END_REF]. In the 1940's a graphite spheroidization treatment was discovered [START_REF] Morrogh | Graphite formation in cast irons and in nickelcarbon and cobalt-carbon alloys[END_REF] where addition of controlled amounts of magnesium, cerium or calcium to the iron melt prior to casting leads to the formation graphite precipitates with a nearly-spherical shape. This new material called "spheroidal graphite cast iron" or "ductile iron" has improved mechanical properties such as ductility and fracture toughness as compared to those of flake-graphite iron and has been used in piping systems, automotive parts and many other applications.

Further research showed that the shape of graphite precipitates in cast irons is very sensitive to small variations in melt composition. This can be beneficial: compacted graphite cast irons, which have a good damping capacity and a good thermal conductivity, can be obtained by reducing the magnesium content of the melt with respect to the amount used in ductile iron production [START_REF] Stefanescu | Solidification and modeling of cast iron-A short history of the defining moments[END_REF][START_REF] Murthy | Characteristics of compacted graphite iron[END_REF][START_REF] Bazdar | Effect of sulfur on graphite aspect ratio and tensile properties in compacted graphite irons[END_REF]. However, this can also be detrimental, as degenerate graphite morphologies can appear in ductile iron castings if certain deleterious elements are present in the melt due to poor charge material quality (purity) or due to the inoculation treatment. The formation of such peculiar microstructures and its important consequences on the mechanical properties of cast irons encouraged intensive research to understand graphite growth mechanisms in cast iron melts. The presence of defects in the graphite lattice, the role of foreign elements and cooling rate are some key points which have been considered to discuss possible graphite growth mechanisms.

There is no doubt that the elongated shape of lamellar graphite results from a preferential growth along the prismatic a axes of graphite. On the contrary, spheroidal growth is associated with an apparent growth along the c axis which is found to be radially orientated in spheroids. Several models have been proposed to describe such a modification of the relative growth rates along a and c axes of graphite. It has been shown that multiple twinning could produce spherulitic crystals in polymer melts [START_REF] Bisault | Spherulitic branching in the crystallization of liquid selenium[END_REF] and a regular branching model following a similar mechanism was proposed to describe spheroidal growth of graphite [START_REF] Faivre | On the mechanisms of spherulitic growth in polymer and iron melts[END_REF]. Some authors have considered that the enhanced c axis growth of graphite was made possible by means of screw dislocations generated by spheroidizing elements adsorbed at growth ledges which would enable spiral growth and eventually lead to the formation of spheres [START_REF] Hillert | The growth of nodular graphite[END_REF]. It has been further stated that spiral growth does not require adsorption of foreign elements in graphite, but would indeed form spherical aggregates consisting of multiple conical helices [START_REF] Double | Cone-helix growth forms of graphite[END_REF][START_REF] Miao | Structure and growth of platelets in graphite spherulites in cast irons[END_REF]. This model had the advantage of reproducing spheroidal structures which are close to those observed experimentally but it has been subjected to further discussion by its own authors [START_REF] Double | The nucleation and growth of graphite-the modification of cast iron[END_REF].

The above models have been criticized since long because there are evidences that graphite growth proceeds mainly along the a axis also in spheroids, even though the overall shape (radial direction of the spheroids) develops along the c direction [START_REF] Loper | Structure of spheroidal graphite in cast iron[END_REF]. For example, some authors have tried to explain the apparent growth along the c axis by a circumferential extension of graphite and/or wrapping of carbon layers [START_REF] Sadocha | The mechanism of graphite spheroid formation in pure FeeCeSi alloys[END_REF]. In this line, it has been previously shown that crystalline defects observed along spheroids' radius could hardly be reconciled with a continuous growth along the c direction [START_REF] Theuwissen | Microstructural characterization of graphite spheroids in ductile iron[END_REF].

It is clear that despite considerable amount of research, the mechanisms of spheroidal graphite growth in metallic melts are still unclear. In particular, if graphite growth in spheroids proceeds along the a direction as does flake graphite, the similarities and possible differences leading to such different overall shapes are worth of further studies. In this article, the microstructural features of different graphite morphologies (spheroidal and flake) appearing in commercial cast irons and synthetic alloys are investigated with microscopy techniques. The results are used to follow through on recent advances in the understanding of this matter [START_REF] Amini | Nucleation and growth kinetics of graphene layers from a molten phase[END_REF] and question some of the models that have been previously proposed.

Materials and experimental methods

Microstructural investigations were first carried out on a commercial spheroidal graphite cast iron kindly provided by Azterlan Metallurgy research center. The composition of this material is reported in Table 1 and details about alloy preparation can be found elsewhere [START_REF] Asenjo | Microstructure investigation of small-section nodular iron castings with chunky graphite[END_REF].

In order to produce flake graphite cast irons, samples of the same spheroidal graphite cast iron were remelted as described previously [START_REF] Lacaze | Distribution and effect of various doping elements during primary graphite growth in cast iron[END_REF]. The samples, placed in graphite crucibles, were heated to 1350 C in a resistance furnace and held at that temperature for 10 min. This stage enabled volatilization of the magnesium added for spheroidization and carbon enrichment of the alloys. From the graphite liquidus in the FeeC phase diagram, the carbon enrichment may be estimated to 0.4 wt.%. They were then cooled in about 10 mine1180 C, i.e. above the FeeC eutectic, and maintained at that temperature for 30 min so as to promote nucleation and growth of graphite crystals in a Fe-rich melt. Following this stage, the samples were rapidly extracted from the furnace and quenched with an air blowing device which led to rapid solidification of the remaining melt at a rate of about 250 C/ min.

Synthetic FeeC alloys were prepared by placing a piece of pure Fe in a graphite crucible (for chemical analysis) and then following the same thermal cycle as above under either air or primary dynamic vacuum (about 5$10 À6 MPa). Similarly, synthetic FeeCeCe alloys were prepared by locating some Ce at the bottom of a graphite crucible and pure Fe on top of it, the amount of Ce being 2% of that of Fe. FeeCeSb alloys were obtained from a Fe-0.5% Sb alloy prepared by induction melting and then following the same procedure than for FeeC samples. According to the FeeC phase diagram, the carbon content of these synthetic alloys should be of the order of 4.6e4.7 wt.% but no chemical analysis was performed because the samples were too small for both metallographic inspection and chemical analysis. One of the FeeCeCe samples investigated in the present work was quenched when the temperature reached 1180 C, i.e. it did not undergo any holding at that temperature.

The samples were cross-sectioned and polished for metallographic examination using a Nikon MA200 inverted metallographic microscope. Specimens for transmission electron microscopy (TEM) were prepared by grinding and polishing these sections to an appropriate thickness. 3 mm discs were then punched (extracted) from the samples, dimpled to about 20 mm and further thinned by ion milling using a Gatan Precision Ion Polishing System (PIPS).

One TEM sample was prepared from the industrial spheroidal graphite cast iron using a JEOL focused ion beam (FIB) following the procedure detailed in Ref. [START_REF] Theuwissen | Nano-scale orientation mapping of graphite in cast irons[END_REF] in order to facilitate homogeneous and localized thinning of a relatively large specimen. All specimens were examined in a JEOL 2100F FEG-TEM operated at 200 kV. Crystal orientation maps obtained on the FIB sample were produced in a JEOL 3010 TEM operated at 300 kV equipped with a Digistar system for automated crystal orientation mapping (ACOM).

Finally, a LEO 435VP scanning electron microscope (SEM) was used to image samples after deep-etching with a 37% HCl-ethanol solution.

Results and discussion

Characterization of graphite in commercial alloys

Spheroidal graphite

Metallographic analysis of the commercial spheroidal graphite cast iron showed that its microstructure consists of spheroidal graphite precipitates of good nodularity (Fig. 1a). An average diameter of 19 mm and a density of 400 mm À2 were determined by image analysis of 4 fields of view of about 0.8 mm 2 on the base material. These precipitates have an irregular outer surface as shown in the micrograph of Fig. 1b obtained after deep etching of the samples. Such a feature has been described as cabbage leaf-like [START_REF] Loper | Structure of spheroidal graphite in cast iron[END_REF] and relates to lateral extension of the outermost graphite layers along the surface of the spheroids, presumably during solidstate cooling of the material.

Previous transmission electron microscopy carried out on spheroidal graphite cast irons revealed that sections of the spheroidal precipitates are made of conical sectors and selected area electron diffraction (SAED) patterns taken in these sectors revealed that c axes of graphite are oriented radially, i.e. roughly parallel to the spheroids' radius in each sector [9,14,19e21]. Misorientations between the (0002) reflections in the diffraction patterns of different locations within one single sector suggest that these fanlike areas are made of domains of similar orientations which have been referred to as crystallites in investigations of carbonaceous materials [START_REF] Mironov | Electron irradiation of nuclear graphite studied by transmission electron microscopy and electron energy loss spectroscopy[END_REF]. These domains are formed during graphite precipitation upon solidification of the iron. As such, these domains of similar orientation can be considered as growth blocks, which, stacked upon each other, form the inner structure of graphite spheroids. These growth blocks have a thickness between 10 and 100 nm [14,19e21] and a length of a few micrometers [START_REF] Monchoux | Morphological changes of graphite spheroids during heat treatment of ductile cast irons[END_REF].

The bright field image (Fig. 2) of one part of a spheroidal graphite precipitate shows an inner structure composed of sectors and a thin outer layer with a less defined structure. Graphite at the periphery of such spheroids appeared to be much less organized than in the bulk of the sectors as shown by spotty rings in the diffraction pattern of Fig. 2a. For comparison, a selected area diffraction pattern was taken in the bulk of a sector with the same aperture size showing one clear orientation (Fig. 2b). It can be concluded from these results that, at the periphery of the spheroids, domains of similar orientation have a smaller size than in the bulk of sectors. Growth blocks are thus smaller in these locations and or present higher relative misorientations between each other. Comparable features have been produced by heat treatments of cast irons and steels [21,23e25]. The authors related the formation of this "microcrystalline graphite" to solid-state carbon deposition during decomposition of the matrix.

A general overview of the inner structure of such spheroidal graphite precipitates was investigated in a previous study by means of Automated Crystal Orientation Mapping in a TEM [START_REF] Theuwissen | Nano-scale orientation mapping of graphite in cast irons[END_REF]. The results, summarized with orientation maps such as those in Fig. 3, clearly show boundaries between neighbouring sectors as well as misorientations within them. Most of the misorientation angles were observed to have values of 10e15 , 20e22 and 27e30 which correspond to known twin angles, though some other values could also be measured. Sharp colour changes within sectors such as the blue to green alternating in the lower right part of Fig. 3 represent rotations of 27 of the c axis of graphite [START_REF] Theuwissen | Nano-scale orientation mapping of graphite in cast irons[END_REF]. This value corresponds to a low energy stacking fault in the graphite lattice as predicted by the coincidence site lattice theory [START_REF] Minkoff | Rotation boundaries and crystal growth in the hexagonal system[END_REF][START_REF] Double | Defects in eutectic flake graphite[END_REF] and determined in recent studies by means of atomistic calculations [START_REF] Zhang | Intrinsic strength and failure behaviors of graphene Grain boundaries[END_REF][START_REF] Zhang | Polycrystalline graphene curved by grain boundary for high performance nanoresonators[END_REF]. According to some authors [START_REF] Double | Cone-helix growth forms of graphite[END_REF][START_REF] Minkoff | Rotation boundaries and crystal growth in the hexagonal system[END_REF], this rotation is likely to occur in a growing graphite crystal and would consequently provide steps for atom attachment [START_REF] Minkoff | Hole formation in crystal growth by surface adsorption of an impurity[END_REF].

After describing the crystal orientation of graphite inside sectors, which consisted mainly in small angle misorientations as well as rotations around the c axis [START_REF] Theuwissen | Nano-scale orientation mapping of graphite in cast irons[END_REF], TEM was used to study the interface between sectors. Diffraction patterns taken over two adjacent sectors (Fig. 4a andb) show a sharp interface clearly defined by a difference in contrast and two distinctive c axis orientations. A high resolution TEM (HRTEM) lattice fringe image taken at the interface between neighbouring sectors in Fig. 4c shows relatively straight fringes on left and right hand sides which are characteristic of a highly graphitic material. Nevertheless, c axes from both parts of the image do not have the same orientation and they are separated by an interface within which the graphene layers are rippled in a transition zone with a width varying from 4 to 9 nm. Indeed, recent studies have shown that grain boundaries in graphene are produced by the introduction of defects which result in an out-of-plane inflexion of the layers [START_REF] Zhang | Polycrystalline graphene curved by grain boundary for high performance nanoresonators[END_REF]. Interestingly, calculations [START_REF] Zhang | Intrinsic strength and failure behaviors of graphene Grain boundaries[END_REF] have shown that the grain boundaries with the lowest formation energy are the most likely to be formed, and these cause inflexion angles of 13.2 , 21.8 and 27.8 close to the ones measured experimentally in this study [START_REF] Theuwissen | Microstructural characterization of graphite spheroids in ductile iron[END_REF].

It thus appears that crystallographic defects such as twins are often present in graphite [31e34], and result in the formation of grain boundaries [START_REF] Zhang | Intrinsic strength and failure behaviors of graphene Grain boundaries[END_REF][START_REF] Zhang | Polycrystalline graphene curved by grain boundary for high performance nanoresonators[END_REF] which accommodate the misorientations between neighbouring sectors.

The observations can be summarized as follows: graphite growth blocks nucleate and grow during solidification of cast iron and are stacked to form conical sectors or fan-like areas. Defects in the graphite structure (misorientations and rotations) are found at the interface between neighbouring sectors creating apparent grain boundaries and also within conical sectors. The spheroidal precipitates formed by multiple sectors can be surrounded by areas presenting small-sized orientation domains which are presumably formed in the solid-state, after solidification.

Flake graphite

Samples of flake graphite irons produced by remelting the initial spheroidal graphite alloy present a microstructure illustrated with Fig. 5a. Relatively large-sized graphite precipitates are seen mainly in the upper part of the samples and at its edges. These are primary precipitates which formed upon cooling from 1350 C and holding at 1180 C. During this stage, graphite precipitates nucleated at the crucible walls and eventually detached to grow freely in the liquid as described by Patterson et al. for similar experiments [START_REF] Patterson | Anwachsversuche an graphit-Impfkristallen aus kohlenstoffges€ attigten Eisenschmelzen[END_REF]. Due to density difference between graphite and iron, these precipitates float to the top of the samples [START_REF] Wetterfal | Solidification process of nodular cast iron[END_REF]. The rest of the samples consists of an iron-rich matrix containing undercooled graphite of smaller size and irregular shape formed during quenching. Such a distribution of graphite precipitates in the sample is inherent to the preparation process and was therefore observed in all laboratory remelted samples which will be presented further. In this study, emphasis will be put on primary graphite. It is generally established that graphite flakes are extended along the a direction and that c axis of graphite is perpendicular to the flake's length. The polarized light micrograph of Fig. 5b shows that the orientation of the c axis is nearly identical in the whole illuminated flake.

Fig. 6 is a bright field TEM image taken at low magnification to show the whole width of a graphite flake. Darker areas on either side of the flake correspond to the surrounding Fe-rich matrix. Selected area diffraction patterns taken at 2 different locations in the flake with a 250 nm aperture show different c axis orientations as seen by the (0002) spots in Fig. 6a andb. [0001] directions in diffraction pattern b are misoriented by 21 with respect to that of diffraction pattern a. These misorientations do not seem to affect the overall growth direction of the flake which remains relatively straight with a fairly homogeneous thickness. However, the presence of such misorientations indicates that the graphite flakes consist of structural entities, or growth blocks which are stacked upon each other to form flakes of a nearly constant thickness. The length of these growth blocks may be much larger than in the case of spheroidal graphite irons, but their thickness is very similar. Comparable to what was observed for spheroidal graphite is the fact that interfaces between growth blocks of different orientations show bent and rippled graphene stacks as shown in Fig. 7. These results demonstrate common features in the inner structures of flake and spheroidal graphite at the microscopic scale.

Graphite growth in synthetic alloys

Synthetic alloys prepared in both air and vacuum atmospheres led to similar microstructures characterized by relatively few primary graphite precipitates as compared to commercial alloys. This is due to the fact that fewer nucleation sites were available in this case because no inoculant was added and the melt was certainly cleaner.

FeeC samples

The optical micrograph in Fig. 8a shows typical flakes observed in a FeeC sample. There are apparent thickness variations along the flakes, indicated by the arrows in the micrograph. This reveals a multi-stage growth in which the steps seen in the optical micrograph of Fig. 8a correspond to growth blocks which extend laterally over the underlying flake as shown in the scanning electron micrograph of a deep etched sample (Fig. 8b). This suggests that lateral extension of consecutive growth blocks results in thickening of the flakes as described by Amini and Abbaschian [START_REF] Amini | Nucleation and growth kinetics of graphene layers from a molten phase[END_REF]. It is interesting to note that no differences were observed between samples processed in air and vacuum.

These observations agree with the views of Minkoff [START_REF] Minkoff | The Physical Metallurgy of Cast Iron[END_REF] who showed that substrate steps, edges or grain boundaries provide nucleation sites for new graphite layers. The thickening rate depends on the nucleation rate of new blocks which then extend along the surface of the flake. Such blocks were reported by Franklin and Stark [START_REF] Franklin | Application of secondary ion mass spectrometry to study of graphite morphology in cast iron[END_REF][START_REF] Franklin | Further use of secondary ion mass spectrometry in the study of graphite morphology control in cast irons[END_REF] who noted a periodic increase of the sulfur signal when carrying out depth profiles of graphite flakes using secondary ion mass spectrometry. The authors attributed this observation to the presence of sulfur atoms on the basal surfaces of graphite, where this element is known to preferentially adsorb. This was considered by the authors as evidence that graphite flakes consist in a buildup of individual lathes.

FeeCeSb samples

FeeCeSb samples also underwent experiments of carbon enrichment in air and vacuum. The role of antimony in graphite growth modification is not yet clear but it is often used in industrial practices to prevent chunky graphite formation in alloys containing rare earths [START_REF] Pan | Effects of Bi and Sb on graphite structure of heavy-section ductile cast iron[END_REF]. The results obtained in this study do not show any significant change in the microstructure with respect to that of pure binary alloys, i.e. the primary graphite precipitates are mainly of the flake-type. Nevertheless, primary graphite flakes appeared to be curved and bent as shown in Fig. 9.

Transmission electron microscopy was performed to examine such precipitates. Fig. 10 shows a long graphite flake, made of several areas of different orientations, delimited by the black dotted lines. The apparent curvature of the flake is caused by multiple changes of orientations in the flake's growth direction, which could have been promoted by the presence of antimony in the alloy. It thus seems that antimony promotes growth defects in the graphite lattice. With the necessary precautions, this could be linked with ab-initio calculations which have shown that Sb reduces graphite twinning energy [START_REF] Bleskov | Effect of antimony on primary graphite growth in cast iron -From ab-initio calculations to experimental observations[END_REF]. As for the case of FeeC alloys, microstructures were similar for samples produced in air and vacuum.

FeeCeCe samples

Samples were prepared in air and vacuum as for the other synthetic alloys. Compared to usual gray irons, the flakes of the samples prepared in air were rather thick and short and could be called platelets as reported in the literature [START_REF] Ruth | Structures de solidification unidirectionnelle des eutectiques fer-carbone[END_REF]. These platelets sometimes show evidence of growth steps (white arrows in Fig. 11a) suggesting that in this case also thickening occurs by a 2D nucleation and growth mechanism.

The samples prepared in vacuum showed peculiar microstructures: primary graphite precipitates with relatively rounded shapes were observed, reminiscent of spheroidal graphite precipitates. For the sample quenched directly after holding at 1180 C, the graphite sections are not complete disks but are made of conical sectors separated by matrix intrusions (Fig. 11b). The extent of these intrusions is variable, but it suggests that several sectors grew independently. It can be referred to as exploded graphite, a term used in the literature to describe a variety of incomplete/degenerate spheroidal graphite precipitates. In some cases, isolated graphite sectors were observed.

The microstructural analysis of the samples reveals that, under similar conditions, removing oxygen (by performing experiments in medium vacuum) enabled the formation of sectors of spheroidal graphite in Ce-bearing alloys. The exploded or spheroidal-like morphology was not achieved without the addition of Ce, even when the FeeC alloys were produced in primary vacuum. This indicates that Ce does first act as an oxygen scavenger as reported in the literature. In this case, it probably gathers the oxygen remaining in the crucible by forming cerium oxides, which is confirmed by the presence of cerium (and oxygen) bearing inclusions found in the samples.

It is worth noting that such exogenous particles can act as nucleants for graphite. The micrograph in Fig. 12, where a graphite particle develops from the outer surface of an inclusion containing Ce and O, is an example of this. Moreover, such proximity between graphite and Ce could ease the incorporation of Ce atoms in graphite [START_REF] Theuwissen | Distribution of some active elements in primary graphite precipitates[END_REF] and possibly lead to the formation of defects in the graphite lattice thus contributing to morphological modifications of the precipitates.

Transmission electron microscopy of a sample that was subjected to the 30 min holding at 1180 C revealed unusual features in the sectors of primary exploded graphite (Fig. 13). Monitoring the c axis orientation over the whole sector showed that in its inner part, [0001] direction is roughly parallel to the radius of the (incomplete) spheroid. The outer part of the sector is a rim with a different contrast in the bright field image, in which the orientation of graphite changes progressively to surround the inner part of the sector. For clarity, approximate c axis orientations are represented schematically in Fig. 13b.

The structural differences of graphite in the two locations suggest that the center and the periphery of the sector grew following different mechanisms occurring at two different growth stages. During cooling from 1350 C to 1180 C, primary graphite nucleates and grows freely in the melt. This growth stage is expected to have led to the formation of the inner part of the sector. During holding at 1180 C, graphite growth is expected to occur in a steady state during which there could be Ce build-up in the liquid, close to the graphite/matrix interface and some of it could be absorbed by the growing graphite.

Energy dispersive X-ray analysis showed the presence of Ce in the outer rim, whereas this element was not detected in the central part of the sector. It is interesting to note that isolated sectors were most often found close to areas which were enriched with cerium (as shown by the numerous Ce-rich particles found at these locations) such as the lower part of the samples, close to where cerium was added prior to melting. This could indicate that the development of complete spheres by lateral extension of graphite sectors may have been hindered by the presence of Ce atoms.

It appears thus that Ce plays a complex role in cast iron melts: it contributes to deoxidizing the melt and would promote nucleation of new growth blocks while limiting their lateral extension. It is therefore likely that at some moment, lateral growth of sectors (isolated or part of an exploded precipitate) would have been blocked and their growth would have proceeded mainly radially, growing outwards along the c axis.

Discussion

The model of 2D nucleation and growth recently described by Amini and Abbaschian [START_REF] Amini | Nucleation and growth kinetics of graphene layers from a molten phase[END_REF] seems thus to govern growth of lamellar graphite in cast irons. Nevertheless, it was described as a layer by layer growth, in which each new layer corresponds to a graphene sheet, i.e., a monoatomic array of carbon atoms. Given the observations of growth blocks reported in this study it is thought that there should be a critical block height or thickness (for given conditions) required for further growth of graphite precipitates instead of atomic layers. This is in line with results from previous studies, in which the assembly of structural units or graphite platelets has been reported to produce a great variety of macroscopic morphologies all presenting the same substructures but following different stacking sequences [START_REF] Yoon | A conceptual model for the structure of catalytically grown carbon nano-fibers[END_REF].

In an earlier study, natural graphite crystals were investigated and showed morphological similarities with those produced in this study [START_REF] Kvasnitsa | Disclinations in unusual graphite crystals from anorthosites of Ukraine[END_REF]. According to these authors, the peculiar shape of the investigated crystals (which was nearly spheroidal) was the result of prominent layered growth, or platelet growth. According to these authors the platelets are likely to be rotated with respect to each other due to defects producing macrosteps and promoting the development of unusual graphite forms by growth of blocks. Bending of graphene sheets, and enhanced c axis growth would be promoted by increasing carbon supersaturation associated with high cooling rates. It is important to point out that these authors reported inclination of the graphite blocks by 9e10 and 13 to the [0001] direction, as was observed in graphite spheroids in this study, and concluded that these misorientations resulted in nearly spherical aggregates. It appears that, opposite to what was stated in the past, screw dislocations are not required to form graphite spheroids and that their contribution to spiral growth around the c axis of graphite is limited [START_REF] Roscoe | Growth of graphite single crystals from iron-carbon solutions[END_REF].

Conclusions

Different types of graphite precipitates found in graphitic cast irons have been investigated. Using mainly transmission electron microscopy, the inner structure of these precipitates was studied and showed that despite morphological differences at the microscopic scale flake and spheroidal graphite precipitates consist of growth blocks stacked upon each other. Crystals grow mainly by a 2D nucleation and growth mechanism, and the final shape of the precipitates is associated with the occurrence of crystallographic defects in the graphite lattice, namely misorientations and rotations during growth. The effect of several elements on graphite morphology was assessed through experiments of primary graphite growth in Fe-rich melts. It is shown that Ce, a known spheroidizing element, does not only play the role of an oxygen scavenger, but might also have a more direct effect on graphite growth, i.e. by limiting lateral growth of graphite. The phenomena described in these Ce-bearing cast irons have also been observed in alloys containing other impurities such as Ca, La, Bi, B and Mg [30,47e49]. These elements, by adsorbing on graphite surfaces and promoting structural defects could lead to a myriad of graphite morphologies, made possible by the high flexibility of the graphite structure under different conditions. The results reported in this study show striking similarities with features of other carbonaceous materials obtained naturally or synthetically in a variety of conditions. Recent advances on graphene research have provided key information to further understand the phenomena occurring during graphite growth and resulting in the inner features of the precipitates observed in this study. All this information seems to point to a general 2-D nucleation and growth mechanism which can be affected by events occurring at the atomic scale level.

Fig. 1 .Fig. 2 .

 12 Fig. 1. a) Optical micrograph of a spheroidal graphite cast iron section and b) SEM image of extracted spheroids. (A color version of this figure can be viewed online.)

Fig. 3 .

 3 Fig.3. Photomontage of TEM images of a diametrical section of graphite spheroid combined with the crystal orientation map of graphite obtained by ACOM; apparent boundaries between sectors are highlighted by white dotted lines (dark areas at the periphery of the nodule are iron-rich particles that were not indexed). (A color version of this figure can be viewed online.)

Fig. 4 .

 4 Fig. 4. a) Bright field TEM image of a boundary between two neighbouring sectors, b) SAED diffraction corresponding to the white circle in the bright field image and c) lattice fringe image of the same interface between neighbouring sectors.

Fig. 5 .

 5 Fig. 5. a) Photomontage of light micrographs showing a sample of flake graphite iron produced by remelting the initial spheroidal graphite alloy and b) polarized light micrograph of graphite flakes. (A color version of this figure can be viewed online.)

Fig. 6 .Fig. 7 .

 67 Fig. 6. Bright field image of flake graphite associated and SAED patterns associated to the labelled areas in the bright field image. (A color version of this figure can be viewed online.)

Fig. 8 .Fig. 9 .Fig. 10 .

 8910 Fig. 8. a) Optical micrograph of typical flakes obtained in FeeC alloy, b) SEM image of a flake after deep etching. (A color version of this figure can be viewed online.)

Fig. 11 .Fig. 12 .

 1112 Fig. 11. Optical micrographs of FeeCeCe alloys a) in air and b) in vacuum. (A color version of this figure can be viewed online.)

Table 1

 1 Composition of the spheroidal graphite cast iron (in wt. %).

	Alloy	C	Si	Mn	P	S	Mg	Cu	Fe
	10F2	3.64	2.05	0.11	0.033	0.015	0.037	0.04	balance

Acknowledgements

The authors would like to thank Muriel Veron for the orientation maps obtained by ACOM, Julien Zollinger for the preparation of the FeeSb alloy and Azterlan for providing some of the materials used in this study.