Edoardo Martino
email: edoardo.martino@epfl.ch

David Santos-Cottin

Florian Le Mardelé

Konstantin Semeniuk

Michele Pizzochero

Kristiāns Čerņevičs

Stefan Klotz

Ludovic Delbes

Benoît Baptiste

Francesco Capitani

Helmuth Berger

Oleg V Yazyev

Ana Akrap
email: ana.akrap@unifr.ch

Structural Phase Transition and Bandgap Control through Mechanical Deformation in Layered Semiconductors 1T-ZrX2 (X = S, Se)

 have been used to grow large single crystals.

High-pressure Raman and infrared spectroscopy: Raman and infrared measurements were performed on the infrared beamline SMIS at the synchrotron SOLEIL. We used a DXR2 Thermo Fisher Raman microspectrometer equipped with a 633nm He-Ne laser as light source, a Peltier cooled CCD detector, a 600 grooves/mm grating, resulting in a spectral resolution of about 4 cm -1 . Infrared measurements were performed on a custom-built horizontal microscope for diamond anvil cells (DACs), equipped with custom Cassegrain objectives, a MCT and a Si diode detectors, for the Mid-IR and the Near-IR respectively [1]. The horizontal microscope was coupled to Thermo Fisher iS50 interformeter with Quartz (NIR) / KBr (MIR) beamsplitters and synchrotron radiation as IR source.

Square samples of about 100 × 100 μm 2 were loaded into a membrane DAC with 400 µm culets.

Stainless-steel gaskets were pre-indented to a thickness of about 50 µm and a hole of diameter 150 um was drilled by electro-erosion. NaCl was used as pressure transmitting medium [2].

Pressure was measured in situ by the standard ruby fluorescence technique. Data was collected every 1 GPa during both compression and decompression, with the last data point collected after the cell was open to release any residual pressure.

Sample thickness for optical transmittance experiments was determined before the measurement by means of optical interference, analyzing the oscillations of the transmitted light intensity as a function of wavelength (Fabry Perot interference). For the computation, the refractive index of the materials at energies below the bandgap is n = 4 [3]. The value of the refractive index is experimentally determined through ellipsometry.

X-ray diffraction:

The high-pressure X-ray diffraction experiments were performed at the Xray diffraction platform of IMPMC at 0.85GPa and 293K, using a membrane diamond anvil cell (DAC) with a culet of 500 μm in diameter. A 250 μm-diameter hole was drilled in a stainless-steel gasket with a thickness of 200 μm pre-indented to 80 μm. The finely ground powders of ZrSe2 and ZrS2 were loaded into the gasket hole together with a ruby ball as pressure sensor [4] (R1 ruby fluorescence method). A 4:1 methanol-ethanol mixture was used as pressure transmitting fluid since it is known to provide hydrostatic conditions to approximately 10 GPa [5]. Before and after each measurement, pressure was determined by recording the ruby fluorescence spectra using a spectrometer with a 532 nm excitation laser. The calculated error for the pressure values is 0.1 GPa and the sensitivity was 0.05GPa.

The DAC was then mounted on a Rigaku MM007HF diffractometer equipped with a Mo rotating anode (λKα1 = 0.709319 Å, λKα2 = 0.713609 Å), VariMax focusing optics and a RAXIS4++ image plate detector. X-ray data were collected at 20 °C. A LaB6 standard sample was measured in the same experimental conditions to calibrate the FIT2D program, the image processing software used to integrate the intensities around the Debye-Scherrer rings and to get the 1D patterns.

Density Functional Theory Calculations: Our first-principles calculations were performed at the density functional theory level as implemented in VASP [6,7]. We adopted the hybrid HSE06 [8] functional for describing the semiconducting systems, and the semilocal PBE functional for the metallic one [9]. Electron-core interactions were described through the projector augmented wave method, while Kohn-Sham wavefunctions were expanded in a plane wave basis set with a cutoff on kinetic energy of 400 eV. The integration over the Brillouin zone was carried out using a 6×6×4 (6×4×2) k-mesh for the low-pressure (high-pressure) crystalline phase of ZrSe2. We relied on the experimental crystal structures determined in our X-ray diffraction experiments.

High-pressure Rietveld refinement of the XRD spectra

The crystal structure refinements were carried out using the Rietveld method as implemented in the Fullprof software [10].

Starting from the known cell parameters of ZrS2 and ZrSe2, the refinements converged rapidly for the known structures at low pressure, considering low amounts of iron (diffraction from the DAC gasket).

For the two structures, preferred orientations were evidenced and their relative parameters refined: (001) for 1T-ZrS2, (-101) for HP-ZrS2, (001) for 1T-ZrSe2, and (010) for HP-ZrSe2.

The peak widths were significantly larger than the instrument resolution (~0.15° 2θ). Starting from an instrumental resolution function, Lorentzian isotropic size (Y) and Lorentzian isotropic strain (X) parameters were refined but only the second one (strain effect) improved the fits. Then, anisotropic refinements of strain parameters were attempted to get better peak profiles. The corresponding spherical harmonic coefficients are given in the following tables but the estimated standard deviations (esd) are too large to discuss the physical meaning of these values.

The low data to parameter ratio at high pressure -mainly due to the coexistence of starting and HP structure and to the peak broadening increasing with the strain -prevent us from a stable refinement of atomic positions. Table 1. Results of the Rietveld refinement from X-ray diffraction analysis (λMo) of ZrS2.

HP-ZrS2

I m3m RBragg = NR Fract. (%) = NR a = NR V = NR

Pressure -Strain relation

Given the experimental X-ray diffraction (XRD) data, recorded during high-pressure experiments, it is possible to compute the change in unit cell parameters of the 1T-phases of ZrS 2 and ZrSe 2 . The data were used to compute the pressurestrain relation for the material.

The engineering strain (in percentage) applied to the material is computed as:

𝜀 𝑎 (𝑃) = 𝑎(𝑃)-𝑎(0) 𝑎(0) • 100 (1)
Where a(0) and a(P) are the same lattice parameter at ambient and pressure P respectively.

Given the hydrostatic conditions of the experiment, only the three diagonal component of the strain tensor are non-zero: εa, εb and εc. Because of the hexagonal lattice symmetry (P-3m1), εa=εb.

(

As a reference, strain is considered positive when tensile and negative when compressive. In the case of our high-pressure experiments, ε is negative as only compressive deformation is produced. In From our results the 1T polytypes appear to have a much higher compressibility than the 2H polytypes. In comparison, at 10 GPa, in-plane strain are εa=3.6% and εa=2.2% for 1T and 2H respectively. For both materials, the out-of-plane compressibility is much higher than the inplane one, as one would expect considering the much weaker van der Waals interaction.

High-pressure Raman scattering

Here we show the Raman spectra at each pressure during compression and decompression. sustain large elastic strains, up to a few percent. However, exfoliable TMDs with highly strain-tunable properties have proven challenging for researchers to identify.

We investigated 1T-ZrS2 and 1T-ZrSe2, exfoliable semiconductors with large bandgaps. Under compressive deformation, both TMDs dramatically change their physical properties. 1T-ZrSe2 undergoes a reversible transformation into an exotic three-dimensional lattice, with a semiconductorto-metal transition. In ZrS2, the irreversible transformation between two different layered structures is accompanied by a sudden 14% bandgap reduction. These results establish that Zr-based TMDs are an optimal strain-tunable platform for spatially textured bandgaps, with a strong potential for novel optoelectronic devices and light harvesting.

When developing semiconductor-based devices for photovoltaic, photocatalytic or photosensing applications, light harvesting efficiencies can be optimised by fine-tuning the size, or even the spatial variation of the electronic bandgap. This can be achieved by mechanical actions, such as strain or hydrostatic pressure. Application of high pressure is an experimentally accessible path for exploring how the properties of a material change upon varying the interatomic distances [1].

Established techniques, such as X-ray diffraction (XRD), can then be used in combination to thoroughly quantify the response of crystalline systems to various mechanical actions, and compute the corresponding compressibility and elastic moduli.

Once a material's mechanical and electronical response to applied forces is known, one can design a device in which, for example, a specific strain induced in the photosensitive component attunes it to a chosen wavelength. A variety of tools and methods for such strain-tuning have been implemented and reported, such as local compression by AFM cantilevers [2], adhesion to an elastomeric substrate [3,4], patterned substrates [5,6], differential thermal expansion between the active material and the substrate [7] and the use of active piezoelectric actuators [8] or microelectromechanical systems [9].

Layered van der Waals materials are especially good candidates for strain-engineered optoelectronic devices [10][11][12], thanks to their proven extremely wide range of elastic strain which they can support (above 5%), when reduced to nanometre thickness by exfoliation [13].

The strain tuning of optoelectronic properties of layered materials, such as transition metal dichalcogenides (TMDs), has already been proposed and successfully demonstrated for the semiconducting compounds containing group 6 metals: 2H-MoS2 and 2H-WSe2 [START_REF] Feng | Strain-engineered artificial atom as a broad-spectrum solar energy funnel[END_REF][START_REF] Xie | Coherent, atomically thin transition-metal dichalcogenide superlattices with engineered strain[END_REF][START_REF] Conley | Bandgap engineering of strained monolayer and bilayer MoS2[END_REF].

Furthermore, tuning the bandgap in TMDs by mechanical action has been used to enhance the performance of semiconducting devices for light harvesting applications [START_REF] De Sanctis | Strain-engineered inverse charge-funnelling in layered semiconductors[END_REF]. The prototypical 2H-MoS2 and 2H-WSe2 are, however, not the only semiconducting layered TMDs. These also include systems with group 4 metals, as Zr and Hf. The corresponding bandgaps range from 1.2 eV to 2 eV, pertinent for optoelectronic and energy conversion applications. These compounds have so far been neglected in comparison to the Mo-and W-based semiconductors, despite their rather promising properties which encompass a high electronic mobility [START_REF] Zhang | Two-dimensional semiconductors with possible high room temperature mobility[END_REF] and a tendency to spontaneously form native oxides with high dielectric constant [START_REF] Lai | HfO 2/HfS 2 hybrid heterostructure fabricated via controllable chemical conversion of two-dimensional HfS 2[END_REF]. The salient differences between group 4 and group 6 TMDs are in the electronic energy bands' orbital character [START_REF] Kolobov | Two-Dimensional Transition-Metal Dichalcogenides[END_REF]. Group 4 TMDs have two fewer valence electrons on the metal atom, and a different metal-chalcogen coordination. Group 6 compounds-for example MoS2-crystallize in the 2H polytype, where the metal is in a trigonal prismatic coordination with six chalcogen atoms. Both conduction and valence bands have a metallic, 𝑑 orbital character, giving the chalcogen a minor role in the electronic structure. In contrast, Zr-and Hf-based TMDs assume the 1T-polytype structure with a six-fold octahedral metal-chalcogen coordination. Their conduction bands originate from the transition metal 𝑑 orbitals, while the chalcogen 𝑝 orbitals are responsible for the valence bands.

Realising a structural phase transition driven by applying a mechanical action is a long-standing desire of the community working on layered semiconductors, such as TMDs. A mechanically induced lattice reconstruction in TMDs can open a venue for strain-controlled semiconductorto-metal conversion [START_REF] Duerloo | Structural phase transitions in two-dimensional Mo-and W-dichalcogenide monolayers[END_REF]. This phenomenon can be utilised in tunable optoelectronic elements, or in clean metal-semiconductor lateral homojunctions [START_REF] Eda | Coherent atomic and electronic heterostructures of single-layer MoS2[END_REF][START_REF] Kappera | Phase-engineered low-resistance contacts for ultrathin MoS2 transistors[END_REF], operated by a stretchable substrate. However for 2H-MoS2 and 2H-WSe2, it is not possible to mechanically induce phase transition that involves in-plane reorganization, like the one that happening in transforming from a 2H into a 1T phases. In fact, structural changes in these compounds cannot be triggered by strain or pressure alone [START_REF] Nayak | Pressure-induced semiconducting to metallic transition in multilayered molybdenum disulphide[END_REF][START_REF] Zhao | Pressure induced metallization with absence of structural transition in layered molybdenum diselenide[END_REF][START_REF] Wang | Pressure-induced iso-structural phase transition and metallization in WSe2[END_REF][START_REF] Caramazza | Effect of pressure on optical properties of the transition metal dichalcogenide MoSe2[END_REF] and require additional high temperatures, laser irradiation or intercalation [START_REF] Guo | Probing the dynamics of the metallic-to-semiconducting structural phase transformation in MoS2 crystals[END_REF][START_REF] Wang | Atomic mechanism of dynamic electrochemical lithiation processes of MoS2 nanosheets[END_REF]. In the present study, we investigated 1T-ZrS 2 and 1T-ZrSe 2 , motivated by the 1T polytypes being considerably more prone to structural instabilities than the 2H polytypes.

In light of the different chalcogen size for the investigated materials, Se being larger than S, their response under applied pressure may be expected to vary. For both mentioned compounds, we observed transitions between two distinct crystalline structures, characterized by different electronic properties. Such lattice reconstructions commonly occur in metallic TMDs, and are responsible for the unusual transport properties of 1T-TaS2 [START_REF] Martino | Preferential out-of-plane conduction and quasi-one-dimensional electronic states in layered van der Waals material 1T-TaS2. npj 2[END_REF], and topological phase transitions in 1T-MoTe2 [START_REF] Wang | MoTe2: a type-II Weyl topological metal[END_REF]. The transition temperatures at which these transformations manifest can be tuned in a broad temperature range by applied pressure [START_REF] Sipos | From Mott state to superconductivity in 1T-TaS2[END_REF][START_REF] Berger | Temperature-driven topological transition in 1T'-MoTe2[END_REF][START_REF] Heikes | Mechanical control of crystal symmetry and superconductivity in Weyl semimetal MoTe2[END_REF]. Indeed, by means of ab initio calculations, Zhai et al. [START_REF] Zhai | Pressure-induced phase transition, metallization and superconductivity in ZrS2[END_REF], predicted a structural phase transition to occur in 1T-ZrS2 at high pressure. To test this experimentally in 1T-ZrS2 and 1T-ZrSe2, we measured their optical and structural properties under quasi-hydrostatic pressure that we applied using a diamond anvil cell.

Figure 1 summarizes a set of structural, infrared and Raman properties for 1T-ZrS2 and 1T-ZrSe2 at ambient pressure. The data were recorded on the same samples as those later used for the high-pressure experiments, and are consistent with the published results [START_REF] Moustafa | Growth and band gap determination of the ZrSx Se2-x single crystal series[END_REF]. Optical transmission measurements (Figure 1e) show transparency in a wide energy range, and the onset of optical absorption at the indirect bandgap energies, which are 1.7 eV for ZrS2 and 1.2 eV for ZrSe 2 . Transparency in the infrared range is accompanied by a remarkably high roomtemperature dc conductivity, measured using a standard 4-point technique, of 3.13 S/cm (0.31 Ωcm) for ZrS2, and 23.81 S/cm (0.04 Ωcm) for ZrSe2. The high intrinsic doping level of these materials [START_REF] Ghafari | Opposite dispersion bands at the Fermi level in ZrSe2[END_REF][START_REF] Moustafa | Growth and band gap determination of the ZrSx Se2-x single crystal series[END_REF], causing their high dc conductivity, is also evidenced by the intense light absorbance at low energies (< 0.5 eV). This is most likely due to in-gap or shallow donor states.

Single crystals of Zr-based TMDs were loaded into a diamond anvil cell for optical measurements under high-pressure conditions. A finely ground powder was used for X-ray diffraction experiments. The applied pressure was monitored using ruby fluorescence [START_REF] Mao | Specific volume measurements of Cu, Mo, Pd, and Ag and calibration of the ruby R1 fluorescence pressure gauge from 0.06 to 1 Mbar[END_REF]. The compressive strain imposed onto the materials at every specific pressure was calculated from our data, and can be found in the SI. The pressure dependence of the Raman and powder XRD patterns are shown in Figure 2. For both materials, the formations of different crystalline structures are evident from the appearance of new peaks in the Raman and XRD spectra. In these specific types of TMDs, but we believe that the different metal-chalcogen coordination, when compared to 2H-polytypes, and the fact that there are two valence electrons less [START_REF] Kolobov | Two-Dimensional Transition-Metal Dichalcogenides[END_REF] on the transition metal, must play a fundamental role.

Using our high-pressure XRD and Raman data, and aided by the work of Zhai et al. [START_REF] Zhai | Pressure-induced phase transition, metallization and superconductivity in ZrS2[END_REF], we succeeded in refining the crystalline structures of HP-ZrS2 and HP-ZrSe2. For ZrS2, mechanical deformation transforms the known layered material into a new layered structure (Fig. 3 a), previously never observed in a TMD. The octahedra are transformed into distorted trigonal prisms. The structure is different from that of a 2H phase, as the triangular bases of the prisms are now parallel to the bc-plane, and their orientation alternates in the a-axis direction. Even more surprising and unusual is the structure assumed by ZrSe2. This initially layered material becomes three-dimensional, a transformation that is reminiscent of the graphite to diamond conversion. The resultant crystalline structure-orthorhombic with the space group Immm-is unusual, and can be observed at ambient conditions only in exotic materials, such as the rareearth dichalcogenide ErSe2 [START_REF] Haase | The Phase Equilibria and Crystal Chemistry of the Rare Earth Group VI Systems. I. Erbium-Selenium[END_REF] and the unconventional heavy-fermion superconductor UTe2 [START_REF] Aoki | Unconventional Superconductivity in Heavy Fermion UTe2[END_REF]. The Zr atoms are coordinated by 8 chalcogens, a configuration that can also be found at ambient pressure in the quasi-one-dimensional crystals ZrS3 and ZrSe3 [START_REF] Furuseth | Crystal Structures of TiS3, ZrS3, ZrSe3, ZrTe3, HfS3 and HfSe3[END_REF].

Rietveld refinement of the powder X-ray diffraction patterns was done for every pressure point (Figures SI. 1 and 2). This was used to determine the possible crystal structure for the highpressure phases, along with the compressibility and volume fraction of the 1T phase as a function of pressure.

Having identified the crystalline structures, the optical and electronic properties of those newly created materials are of immediate interest. By measuring optical transmittance as a function of pressure, we investigated the presence and magnitude of the optical bandgap (further details can be found in the SI). For ZrS2, at low pressures (0 to 2 GPa) the bandgap of the 1T phase is reduced, like it is commonly seen in other TMDs [START_REF] Nayak | Pressure-induced semiconducting to metallic transition in multilayered molybdenum disulphide[END_REF][START_REF] Zhao | Pressure induced metallization with absence of structural transition in layered molybdenum diselenide[END_REF]. Above 3 GPa, as the new phase appears, it is possible to identify the contribution to absorption from a second bandgap, which is approximately 200 meV smaller than that of the native 1T phase. As pressure further increases, the bandgap decreases, offering additional control over the optical properties. At the same time, the volume ratio between the two phases changes, as seen from the XRD data (SI), further modifying the optical spectrum. To confirm our estimate of the optical bandgap, we performed first-principle calculations on the ZrS2 HP phase at 6 GPa. The computed band structure has an optical bandgap of 1 eV, in perfect agreement with the value extracted from the experimental data (Figure SI 8). The origin of a small suppression of optical transmittance in the 0.5 -1 eV energy range remains unclear.

ZrSe2 changes its properties more dramatically. We observed a continuous reduction of the bandgap of the 1T phase up to 5 GPa,. As the new crystalline phase appears, the optical transparency is suppressed for all frequencies. This means that HP-ZrSe2 is metallic [START_REF] Peter | Fundamentals of semiconductors: physics and materials properties[END_REF], or possibly a narrow-gap semiconductor with a bandgap of less than 0.1 eV. The evolution of the optical bandgaps extracted from the optical transmittance data is summarized in Figure 3 c,f.

In order to further address the dramatic pressure-induced semiconductor-to-metal transition observed in ZrSe2, the band structure for the proposed high-pressure phase was computed by means of semilocal Perdew-Burke-Ernzerhof functional. Specifically, we determined the electronic band structures of both the low-pressure (0 GPa) and the high-pressure (9.5 GPa)

phases by means of density functional theory. The low-pressure phase is found to be a semiconductor, with an indirect bandgap of 0.9 eV (see Figure 4a). This value is in agreement with our optical measurements shown in Figure 3e and 3f, and perfectly matches the results of recent ARPES experiments [START_REF] Ghafari | Opposite dispersion bands at the Fermi level in ZrSe2[END_REF]. We found that the valence band maximum is located at the centre of the Brillouin zone, while the conduction band minimum occurs at the high-symmetry L point. As far as the high-pressure phase of ZrSe2 is concerned, the present optical measurements cannot distinguish whether this system is a metal, or a semiconductor with a bandgap below 0.1 eV, which is the lowest energy we could probe with our experimental setup.

Our calculations reported in Figure 4b clearly indicate a metallic character, with a number of bands crossing the Fermi level. Overall, our calculations paint a picture in accordance with the experimental observations.

To conclude, we have experimentally demonstrated room-temperature pressure-driven structural phase transitions in the semiconducting layered transition metal dichalcogenides 1T-ZrS2 and 1T-ZrSe2. The two materials are isostructural and isoelectronic, but undergo completely different lattice reconstructions, and their physical properties evolve in distinct manners. Under applied pressure of 3 GPa, ZrS2 is transformed irreversibly into a new layered structure, with a consequent reduction of the bandgap by 200 meV. As for ZrSe2, once pressure is above 8 GPa, the layers merge into a 3D structure. This collapse results in a semiconductorto-metal transition, as confirmed by ab initio calculations. In addition, the bandgap of the starting 1T phase reduces under compression at the rate of 50 meV/GPa, giving a further tunability over the materials' optical properties. Our demonstration of structural transformations induced by mechanical actions, with the associated dramatic changes in their electronic properties, grants the long-awaited additional functionality in TMDs, enabling us to control structure and properties on-demand. New opportunities in the field of strain-designed devices-charge funnelling and clean metal-semiconductor homojunctions in optoelectronic components-are now within reach. Further research needs to be done in the direction of understanding the dynamics of the lattice reconstruction, in particular for ZrSe2 because of its close reminiscence to the graphite to diamond transformation. TOC figure

  .25 (0), S 004 =5.67(4.86), S 112 =0.23(11.30)

a 3 Strain 3 Strain 3 Pressure 11 a = 3 .

 333113 = b = 3.6473 (9) Å c = 5.7208 (1) Å V = 65.907 (24) Å 3 Strain Parameters (for profile only, large esd) Parameters (for profile only, large esd) S 400 = 6.98(7.71), S 040 = 21.10 (31.36), S 004 = 1.43 (1.44), S 220 = 10.68 (72.44), S 202 =0, S 022 =0 Parameters (for profile only, large esd) S400= 0.48 (41), S040= 1.03 (1.55), S004= 0.09 (7), S220= 1.55 (3.81), S202= 0, S022=0 Fe I m3m RBragg = 0.5 % Fract. (%) = 0 (1) a = 2.7413 (497) Å V = 20.599 (647) Å 5894 (239) Å b = 5.4119 (111) Å c = 12.0219 (116) Å V = 233.527 (552) Å 3Strain Parameters (for profile only, large esd) S400= 1.09(0.74), S040= 2.03 (2.72), S004= 0.07 (0.08), S220= NR, S202= 0.67 (0.68), S022=0

  Figure SI 1, we report the computed values, extracted from high-pressure powder XRD for 1T-ZrS2 and 1T-ZrSe2. As a comparison, Figure SI 2 shows the compressibility as function of pressure for the semiconducting 2H polytypes (MoS2, WS2 and WSe2), computed from published results[11,12,13].

Figure SI 3

 3 Figure SI 3. a) 1T-ZrX2 unit cell. The arrows point in the direction of positive strains (tensile) along 2 principal directions, εa is in-plane and εc is out-of-plane. b) Experimentally determined pressure-strain relation for 1T-ZrS2. c) Experimentally determined pressurestrain relation for 1T-ZrSe2.

Figure SI 4

 4 Figure SI 4. a) 2H-MeX2 unit cell. The arrows point in the direction of positive strains (tensile) along 2 principal directions. b) Experimentally determined pressure-strain relation for 2H-MoS2, 2H-WS2 and 2H-WSe2.

Figure SI 5 .

 5 Figure SI 5. Raman spectra of ZrS2 collected during compression and decompression. All curves are shifted vertically by an equal amount.

Figure SI 6 .

 6 Figure SI 6. Raman spectra of ZrSe2 collected during compression and decompression. All curves are shifted vertically by an equal amount.

Figure SI 7 .Figure SI 8 .

 78 Figure SI 7. Representative extrapolation of the indirect band gap for the high-pressure phase of ZrS2 (HP-ZrS2) at different pressures.

ZrS2, the new

  crystalline phase starts to appear at 3 GPa, which corresponds to 1 % in-plane compression and 4% out-of-plane compression(Figure SI. 3). With increasing pressure, a larger fraction of the sample is converted into this new phase, with the low-and high-pressure phases coexisting over a broad range of pressures. The transformation into the newly created phase, here called HP-ZrS2, is irreversible upon full decompression to 1 atm. ZrSe2 remains structurally stable up to 8 GPa, equivalent to 3 % in-plane compression and 7% out-of-plane compression (Figure SI. 3). Upon further increase of pressure, a structural transformation takes place, with a more rapid conversion of the material compared to the S-based compound. The high-pressure phase (HP-ZrSe2) persists down to 2 GPa when pressure is released. Below 2 GPa, the initial structure is recovered, making the transformation reversible, but with a wide hysteresis. The progression and reversibility of the phase transitions as a function of pressure are schematically depicted in Figure2g,h. The different behaviour between the two isostructural materials-different transformation onset pressures as well as the presence or absence of reversibility-suggests a possible difference in the crystalline structure of the highpressure phases. It is not yet clear what makes this structural transformation favourable only for

Figure 1 .

 1 Figure 1. Structural and optical properties of 1T-ZrS2 and 1T-ZrSe2. a) Crystal structure of both ZrS2 and ZrSe2. Red and yellow spheres represent the metal and chalcogen atoms, respectively. The octahedra formed by the chalcogens around the central Zr atom are shown. b) Photographs of the single crystals used for the experiments. The bright red colour of ZrS2 is the result of its optical bandgap in the visible range. c,d) Rietveld-refined powder X-ray diffraction (XRD) of 1T-ZrS2 and 1T-ZrSe2 measured at low pressure (< 1 GPa) inside the diamond anvil cell. Red dots are the experimental data, and the black lines are the simulated patterns for the known crystal structures. The diffraction pattern from the steel gasket is also included into the refinement. e) Optical transmittance through thin samples (6 μm and 2 μm for ZrS 2 and ZrSe 2 , respectively), where the arrows indicate the positions of the indirect bandgaps at the onset of optical absorption. f) Raman spectra with the indexed A1g and Eg phonon modes, shared by both materials due to their identical crystalline symmetry. Raman active modes are: in ZrS2, A1g = 334 cm -1 and Eg = 249 cm -1 ; in ZrSe2, A1g = 195 cm -1 and Eg = 146 cm -1 .

Figure 2 .

 2 Figure 2. Evidence of pressure-induced structural phase transitions a,b) Raman spectra of ZrS2 during compression (a) and decompression (b). The dotted white and purple lines are guides for the eyes in the low-and high-pressure phases, respectively. The compression curve shows the onset of transformation at 3 GPa. During decompression the high-pressure phase remains stable down to 1 atm. c,d) Pressure evolution of the Raman spectra of ZrSe2. Data during compression (c) shows the appearance of a new crystalline phase at 8 GPa, and the initial phase is restored during decompression below 2 GPa (d). e,f) Powder X-ray diffraction at different pressures for ZrS2 (e) and ZrSe2 (f), the purple lines mark the diffraction peaks originating from the new crystalline phase. g,h) Schematic representation of the pressure driven transformation. The blue and purple regions represent the ambient-pressure and high-pressure phases, respectively. The red crosses show the volume fraction of the 1T phase as a function of pressure during compression.

Figure 3 .

 3 Figure 3. High-pressure structures and switching of optical properties a) Crystal structure of HP-ZrS2. The layered structure is evident, while the original octahedral coordination is distorted into an unusual trigonal prismatic configuration. b) Optical transmittance as function of pressure for ZrS2. c) ZrS2 bandgap evolution under pressure d) Crystal structure of HP-ZrSe2. The atomic lattice is three-dimensional, a rare configuration that exists at ambient pressure in ErSe2 and UTe2. e) Optical transmittance measured up to 8 GPa. As the new phase appears, optical transmittance becomes zero due to its metallic properties. f) ZrSe2 bandgap evolution under pressure. The grey area in the panels b and e are the 0.23-0.28 eV region obscured by diamond absorption.

Figure 4 .

 4 Figure 4. First-principles calculations of the electronic structure of ZrSe2. Electronic band structure of the (a) low-pressure and (b) high-pressure crystalline phases of ZrSe2, together with the high-symmetry points in the corresponding first Brillouin zones.

3 Strain Parameters (for profile only, large esd)

	S400=45 (83), S040=55(66), S004=140(306), S220=138(127), S202=0
	S022=0, S121=-0, S301=0, S103=0
	Fe	I m3m	RBragg = 15.0 %	Fract. (%) = 6 (1)
	a = 2.8616 (30) Å
	V = 23.433 (42) Å 3
		Pressure 6GPa
	1T-ZrS2	P -3m1	RBragg = 6.2 %	Fract. (%) = 37 (4)
	a = b = 3.5904 (17) Å	c = 5.4493 (52) Å
	V = 60.837 (71) Å 3

P 21/m RBragg = 17.8 % Fract. (%) = 38 (8) a = 6.4811 (104) Å b = 3.7610 (10) Å c = 5.0823 (126) Å = 108.922(50) deg V = 117.188 (341) Å

Strain Parameters (for profile only, large esd)

	S400=26 (29), S004=0 (8), S112=0
	HP-ZrS2	P 21/m	RBragg = 11.6 %	Fract. (%) = 55 (6)
	a = 6.4766 (63) Å	b = 3.7235 (25) Å	c = 4.9906 (75) Å	= 108.549(54) deg
	V = 114.100 (218) Å 3

Strain Parameters (for profile only, large esd)

	S400=45 (28), S040=146(52), S004=46(60), S220=88(90), S202=124 (104)
	S022=0, S121=-0, S301=0, S103=0
	Fe	I m3m	RBragg = 7.1 %	Fract. (%) = 8 (1)
	a = 2.8587 (34) Å
	V = 23.363 (48) Å 3
		Pressure 8GPa
	1T-ZrS2	P -3m1	RBragg = 8.7 %	Fract. (%) = 28 (6)
	a = b = 3.5789 (28) Å	c = 5.4162 (99) Å
	V = 60.081 (128) Å 3

Strain Parameters (for profile only, large esd)

	S 400 =109 (85), S 004 =63 (69), S 112 =126 (138)
	HP-ZrS2	P 2 1 /m	R Bragg = 15.0 %	Fract. (%) = 62 (8)
	a = 6.4791 (80) Å	b = 3.7064 (41) Å	c = 5.0004 (123) Å	= 108.88(11) deg
	V = 113.620 (337) Å 3

Strain Parameters (for profile only, large esd)

	S400=64 (26), S040=169 (65), S004=262 (126), S220=162 (109), S202=0
	S022=0, S121=-0, S301=0, S103=0
	Fe	I m3m	RBragg = 10.1 %	Fract. (%) = 10 (1)
	a = 2.8641 (29) Å
	V = 23.493 (41) Å 3
		Pressure 10GPa
	1T-ZrS2	P -3m1	RBragg = 12.3 %	Fract. (%) = 28 (11)
	a = b = 3.5563 (14) Å	c = 5.3731 (90) Å
	V = 58.850 (104) Å 3

Strain Parameters (for profile only, large esd)

	S400=155 (172), S004= 383 (299), S112= 0
	HP-ZrS2	P 21/m	RBragg = 15.6 %	Fract. (%) = 65 (9)
	a = 6.4673 (94) Å	b = 3.6731 (21) Å	c = 4.9983 (138) Å	= 109.23(13) deg
	V = 112.106 (355) Å 3

Strain Parameters (for profile only, large esd)

  Results of the Rietveld refinement from X-ray diffraction analysis (λ Mo ) of ZrS 2 .

	S400=11 (22), S040=200 (64), S004=125 (118), S220=210 (85), S202= 286 (134)
	S022=0, S121=-0, S301=0, S103=0
	Fe	I m3m	RBragg = 8.9 %	Fract. (%) = 7.1 (2)
	a = 2.8558 (31) Å
	V = 23.290 (44) Å 3
	Figure SI 1

Table X .

 X Results of the Rietveld refinement from X-ray diffraction analysis (λMo) of ZrSe2.

		Pressure 0.4GPa
	1T-ZrSe2	P -3m1	RBragg = 7.72 %	Fract. (%) = 98 (8)
	a = b = 3.7656 (2) Å	c = 6.0953 (30) Å
	V = 74.854 (37) Å 3

Strain Parameters (for profile only, large esd)

	S400=1.16 (34), S004=5.76 (4.31), S112=3.32 (4.50)
	Fe	I m3m	RBragg = 6.65 %	Fract. (%) = 2 (2)
	a = 2.8756 (8) Å
	V = 23.779 (11) Å 3
		Pressure 1.7GPa
	1T-ZrSe2	P -3m1	RBragg = 8.28 %	Fract. (%) = 98

Acknowledgements

The authors acknowledge illuminating discussions with Laszlo Forrό, Nathaniel Miller, Anna Celeste, Ferenc Borodincs, Quansheng Wu, Zoran Rukelj and Alexey Kuzmenko. We would like to thank Philippe Rosier and Nicolas Dumesnil (from IMPMC) for the production of mechanical parts which were useful for mounting the diamond anvil cell on the diffractometer and high-pressure X-ray diffraction experiments. E.M. and K.S. acknowledges funding from the Swiss National Science Foundation through its SINERGIA network MPBH and grant No. 200021_175836. A.A. acknowledges funding from the Swiss National Science Foundation trough project PP00P2_170544. M.P., K.C. and O.V.Y. are financially supported by the Swiss National Science Foundation through the grant 172543. First-principles calculations were performed at the Swiss National Supercomputing Centre (CSCS) under the project s832. Experiment at Synchrotron Soleil was supported by the grant proposal ID 20190927.

Data availability

The data that support the findings of this study are available from the authors (E.M. and A.A.) upon reasonable request.

Supporting information

The Supporting Information is available free of charge at Experimental details: Crystal synthesis, High-Pressure Raman and infrared spectroscopy, X-ray diffraction, density functional theory calculations.

Competing interests

The authors declare no conflict of interest.


Abderrahmane Hamdi

Ioana Silvia Hosu

Yannick Coffinier
email: yannick.coffinier@univ-lille.fr

Influence of buried oxide layer of nanostructured SOI surfaces on matrix-free LDI-MS performances

Keywords: nanostructured SOI, BOX thicknesses, metal assisted chemical etching, massspectrometry, laser desorption ionization, peptides

In this paper, we report on the nanostructuration of the silicon crystalline top layer of different "home-made" SOI substrates presenting various buried oxide (BOx) layer thicknesses. The nanostructuration was achieved via one-step metal assisted chemical etching (MACE) procedure. The etched N-SOI substrates surfaces were then characterized by AFM, SEM and photoluminescence. To investigate their laser desorption/ionization mass spectrometry performances, the different surfaces have been assessed towards peptide mixtures. We have shown that the matrix-free LDI process occurred from surface heating after laser irradiation and was fostered by thermal confinement in the thin nanostructured Si surface layer. This thermal confinement was enhanced with the increase of buried oxide layer thickness till an optimal thickness of 200 nm for which the best results in terms of signal intensities, peptides discrimination and spot to spot and surface to surface variations were found.

Introduction:

Matrix-free laser desorption/ionization (LDI) process known as SALDI (surface-assisted laser desorption/ionization) has been developed to limit the background coming from adducts and fragments of organic matrices present in the low molecular mass region during matrix assisted laser desorption ionization mass spectrometry (MALDI-MS). The detection of small compounds, below 700 Da, is thus hampered when MALDI-MS is performed. SALDI process mainly uses nanomaterials made of semi-conductors, metals, metal oxides and carbon-based structures. They are presenting several advantages for the analysis of small compounds 1 : i) no need of co-crystallization of the analytes with organic matrix, thus suppressing the intense background from matrix peaks [2][3][4] . ii) simplicity of sample deposition and iii) better reproducibility of results. All these advantages make SALDI surface suitable for high throughput analyses and make possible their integration in lab-on-chip devices 5,6 .

Moreover, by controlling their shape, morphologies, porosities and dimensions (particles, nanowires…) their optical, electrical and thermal properties can be tuned for SALDI-MS applications 1,7 . From decades, SALDI surfaces presenting different shape and dimensions such as 0D (particles), 1D (nanowire/nanotubes) or 2D films (graphene, MoS 2 ) have been tested as LDI materials [8][9][10][11] . Recently, a new class of plasmonic metal materials has emerged as LDI-MS substrates. Indeed, gold particles, mesoporous organosilica@Ag janus particles or magnetic@SiO 2 @Ag particles have shown very promising applications for the detection of nutrients, cervical and lung cancers and bacterial metabolism analysis. [12][13][14][15][16][17] However, porous silicon is still one of the most SALDI materials used, notably due to its high specific surface area, straightforward preparation, UV absorption, antireflective and low thermal properties (heat confinement) (see table 1) 18 .

The matrix-free LDI process is mainly based on thermal driven mechanism. It comprises several occurring concomitant phenomena such as energy absorption (from laser), desorption and ionization steps 19 . So, to perform SALDI-MS analyses, the interfaces should be a good photon absorbent at =355 nm (laser wavelength with photon energy value of 3.49eV) 19 . Then, via non-radiative recombination of electrons and holes (heat dissipation via phonons), the desorption/ionization of analytes can take place. It's well known that surfaces morphology, dimension, porosity, density and periodicity are sturdily related to their physical properties such as optical absorption and thermal dissipation 20 . By tuning the nanostructures' dimension, a better control of the LDI performances is expected. Kim et al., have recently shown that commercial SOI substrate can be used as SALDI surface for the MS detection of insulin 21 . In this study, authors have demonstrated that the laser energy is trapped into the thin top active silicon layer supported by a thermal insulator material (SiO 2 ). This buried oxide (BOx) presents a lower thermal conductivity ~ 1.3 Wm -1 K -1 vs ~40-60 Wm -1 K -1 for top silicon layer (thicknesses of 50-100nm) (following the Fuchs-Sondheimer model) that is to be compared to thermal conductivity value of 150 Wm -1 K -1 for bulk Si 22 . The presence of underlying insulator material may provide an effective heat diffusion barrier which enables the energy to stay confined within the thin surface layer with low thermal conductivity, promoting heat confinement. In addition, it's well known that the porosification drastically decreased the thermal conductivity of silicon nanostructures 23,24 .

So here, what we propose is to take advantage of the nanostructuration of the silicon layer of SOI substrate via metal-assisted chemical etching (MACE) method to obtain a low thermal conductive surface (N-SOI) for LDI-MS detection of small compounds. By doing so, lower laser fluence to perform LDI process is expected. As model compounds, we choose calibration peptides and antibiotics.

II) Materials & Methods a) Materials

Silicon wafers have been purchased from Sil'tronix Silicon Technologies (Archamps, France). All cleaning and etching reagents were of VLSI grade. Sulphuric acid, 96% (H 2 SO 4 ) from Technic (France), hydrogen peroxide 30% (H 2 O 2 ) and nitric acid 65% (HNO 3 ) from Carlo Erba, hydrofluoric acid 50% (HF) and ammonium fluoride (NH 4 F) were supplied by BASF. Silver nitrate (AgNO 3 ), acetone, hexane, dichloromethane (CH 2 Cl 2 ), trifluoroacetic acid (TFA), ammonium citrate dibasic, methanol (MeOH), isopropyl alcohol (IPA) and ethanol (EtOH) were purchased from Sigma-Aldrich. Otcadecyltrichlorosilane (OTS) was purchased from ABCR (Germany). Peptide mixtures were obtained from Applied biosystem. An ultrapure water (Milli-Q, 18 MΩcm) system was used for solutions preparation and rinsing steps.

b) Nanostructures fabrication

The nanostructuration of the c-Si top layer of the home-made SOI surface was achieved by chemical etching in NH 4 F/HNO 3 /AgNO 3 aqueous solution 25,26 . Home-made SOI substrates were fabricated as follow: The silicon surface ((100), p-type, low resistivity) was cleaned in subsequent baths of acetone and isopropyl alcohol, followed by a rinsing step in MilliQ water. Then, the surface was plunged in a piranha solution (H 2 SO 4 /30% H 2 O 2 ; 1/1 (v/v)) during 20 min at RT, followed by copious rinsing with Milli-Q water. The clean silicon surfaces were then thermally oxidized in a Low Pressure Chemical Vapor Deposition (LPCVD) furnace (Tempress) by using a well calibrated process. The thicknesses of the oxidized silicon were measured by ellipsometry (Horiba Jobin Yvon). 5 thicknesses were obtained with the following values: 0, 50, 100, 200 and 400 nm of SiOx. Then, 150 nm of a crystalline silicon layer was deposited by Low Pressure Chemical Vapor Deposition. The thickness was measured by ellipsometry and confirmed by SEM cross-view measurement. For the nanostructuration of SOI surfaces, they were dipped in NH 4 F/HNO 3 /AgNO 3 (6.00 M/5.73 M/0.02 M) aqueous solution at room temperature during various times: from 10 s to 5 min, then followed by a copious rinsing step with deionized water. To ensure the complete removal of silver dendrites formed during the chemical etching, the surfaces were immerged in HNO 3 (65%) during 10 min. Finally, the nanostructured SOI (N-SOI) surfaces were copiously rinsed with deionized water and dried under a stream of nitrogen. It has to be noticed that the same process was applied on crystalline Si wafer noted N-cSi. SOI (unstructured) and N-cSi surfaces were used as references.

Safety considerations

H

c) Silanization of surface by octadecyltrichlorosilane (OTS)

All surfaces were submitted to UV/ozone treatment (UV-O Cleaner, Jelight Company, Inc., 4mW/cm 2 at 220 nm) during 15 min to generate silanols groups (activation). Then, the surfaces were plunged in 10 -3 M solution of OTS in hexane during 4h at room temperature in a dedicated glovebox [25][26][27] . The surfaces were then rinsed with CH 2 Cl 2 , isopropyl alcohol and dried by a stream of nitrogen.

d) Characterization of interfaces

Scanning electron microscopy (SEM)

SEM images were obtained using an ULTRA 55 (Zeiss) electron microscope equipped with EsB, high efficiency In-lens SE and Everhart-Thornley Secondary Electron detectors and an energy dispersive Xray analysis device (Brüker) for EDX measurements.

Atomic force microscopy

Atomic Force microscopy (AFM) imaging was achieved in the tapping mode with a Dimension 3100 Model AFM (Veeco, Santa Barbara, CA) with a Nanoscope IV controller from Digital Instruments under ambient conditions. Silicon cantilevers (Nanoworld) with spring constants of 42 Nm -1 and resonant frequencies between 250 and 400 kHz were used.

Photoluminescence measurement

Photoluminiscence spectra have been recorded at 351.1 nm (argon laser (P=10 mW)) at room temperature by using an IHR 320 spectrometer from HORIBA Jobin Yvon. A 10× microscope objective was used for both producing a beam waist of 0.5 µm and collecting the backscattered PL light from samples. A 360 nm emission cutoff filter was used during measurements.

Reflectivity measurement

An UV-vis spectrophotometer from Perkin-Elmer (with an integrating sphere) was used for measuring sample reflectance. Spectra were registered between 200 and 800 nm at an incident light angle of 45°.

Contact angle measurements

A goniometer system DIGIDROP (GBX, France) was used for measuring the water contact angles. All measurements were performed by using a water droplet of 1µL. All measurements have been done in clean-room facilities environment. All presented values are average of 4 measurements done on different surface locations.

e) LDI-MS analysis

An Ultraflex TOF/TOF instrument (Bruker Daltonics, Wissembourg, France) with a pulsed Nd:YAG laser (=355 nm at 100 Hz) was used. All mass spectra were acquired in the positive and reflectron modes. The laser fluence was adjusted for each experiment and ions were detected from m/z 800 to 6500. The S/N threshold was fixed at 3 for better reliability. Data acquisition and processing were achieved with Flex Control and Flex Analysis softwars, respectively. Instrument calibration was achieved prior each experiment. N-SOI interfaces were fixed on stainless-steel MTP TLC adapter (Bruker) with a conductive tape and directly introduced into the mass spectrometer ion source. All experiments were repeated at least three times to ensure better reproducibility.

d) Peptide solutions

The peptide solutions were composed by: i ). All peptides were resuspended in water/acetonitrile 50/50 at 25 pmol/µL in presence of 0,3% TFA. Before each SALDI-MS experiment, peptide solutions were diluted in ammonium citrate (1mM) at the desired concentration (5pmol, 500 fmol and 50 fmol). Then, 0.5µl was deposited on N-SOI surface and let dried at room temperature under ambient atmosphere.

III-Results and discussion

Morphology of the SALDI surface is crucial in the matrix-free LDI processes. Indeed, it is known that interactions of nanostructured materials (dimensions <500 nm) with laser irradiation lead to more efficient desorption and ionization of compounds of interest 20 . This can be explained by the fact that for such nanostructures' dimensions, optical absorption and thermal dissipation are enhanced 20 .

From this, we took advantages of the SOI substrate features i.e. thin layer of crystalline Si layer supported on buried silicon oxide layer and its nanoengineering to finely control the LDI-MS performances.

In this study, we set up "home-made" SOI surface composed by an active silicon top layer of 150 nm with various thicknesses of BOx layer. So, 5 types of SOI were chosen showing various BOx thicknesses i.e. 0, ~50, ~100, ~200 and ~400 nm (Figures S1 &S2). Then, the fabrication of the nanostructured layer was achieved by submitting SOI substrates to MACE method as schematically represented on figure 1. In figure 2 is represented SEM images of SOI substrate (BOx~200 nm) before and after etching. After 10 seconds of etching, we can clearly see the nanostructures formation (slightly white by eye) with no decrease of the polysilicon or SiOx layers. Interestingly, for etching time greater than 10 seconds, the Si dissolution seems to take place laterally (Figure S3). Indeed, longer times of etching led to bigger holes formation, until the total disappearance of the top silicon layer of SOI (5 min of etching) (Figure S3). In figure S4 presented SEM images of crystalline Si wafer submitted to MACE process, noted N-cSi. AFM characterization gave rough mean square (RMS) values of 0.1 and 1.7 nm for SOI and N-SOI, respectively (Figure 3), thus showing the obvious increase of surface roughness after the chemical etching.

III-2 Optical characterizations

Anti-reflective property of our N-SOI surfaces is crucial since it will help them to maximize photons absorption from the laser irradiation. The total reflectivity spectra for both bare SOI and N-SOI surfaces have been recorded from wavelengths ranging from 250 to 800 nm (Figure S5). Both SOI (BOx=0nm) and N-SOI (BOx=0nm) surfaces present similar reflection spectra profile with lower reflectance for the etched one due to lower refractive index (Figure S5) 28 . For other SOI surfaces with BOx=50, 100, 200 and 400 nm, reflection spectra are quite different and are divided in two zones, a metallic-type reflection at small wavelengths <425 nm and Fabry-Perot cavity type reflection (fringes) at lower wavelengths, typical of SOI substrate and mainly due to reflections at the Si/SiO 2 interfaces (Figures 4 andS5) 29 . After etching, a reflectance reduction (increase light absorption) is observed that can be attributed to the strong light scattering and diffraction occurring at the nanostructured interfaces (Figures 4 andS5). To summarize, by achieving chemical etching of SOI surface, we increased the surface porosity, leading to better anti-reflective properties (30% of reflectivity for N-SOI vs 45% for SOI) at 355 nm (wavelength used for LDI-MS) for all considered BOx thicknesses. This better anti-reflective property makes N-SOI better photon absorber than bare SOI surface thus suggesting a better energy transfer that can take place, from N-SOI surface to analytes during LDI process 30 .

Moreover, the laser penetration depth in the silicon nanostructures at the considered wavelength (=355 nm) is lower than 100 nm, by taking into account the nanostructure optical absorption and reflection in this energy range. As the thin nanostructured active silicon layer is ~150 nm, we can expect an exclusive photon absorption by the top layer during pulsed laser irradiation. In addition, in the case of SALDI surfaces, analytes have to be uniformly distributed within the nanostructured layer to be efficiently desorbed. Indeed, it was noticed that when the thickness of a porous silicon layer is too high, analytes could be trapped deeply inside the nanostructured layer 27,31 . Then, albeit the energy was accordingly transferred to analytes, the desorption phenomenon could be less efficient with as main consequence, a decrease of the detection sensitivity. In our case, analytes deposited on N-SOI will be well distributed and won't have the possibility to go deeper because of the underlying SiO x . Finally, the presence of SiO x underneath the nanostructured silicon layer will enhance the heat confinement inside the nanostructured silicon layer after laser irradiation, improving the efficiency of the LDI-MS process.

The photoluminescence measurements were also performed of both surfaces and spectra are displayed in figure 5. The excitation wavelength was fixed at 351 nm. We can observe a broad PL spectrum from 450 nm to 900 nm for both substrates. By comparing both intensities, no increase of PL was observed, meaning that no quantum confinement occurred after etching. However, a centered peak at 625 nm for the N-SOI substrate (BOx 200 nm, blue curve) which is blue-shifted (~675 nm) compared to PL obtained from SOI substrate (BOx 200 nm, black curve). This is a further indication that the PL originates from surface defects and from the silicon oxide layer 32 . We can then postulate that the adsorbed energy will be mainly dissipated by non-radiative mechanism, leading to heat generation 33 .

III-3 Chemical modification & LDI-MS applications III-4-1 Water contact angle (WCA) measurements

The freshly prepared N-SOI surfaces present superhydrophilic behaviors with water static contact angle (WCA) values of ~0°, that is mainly due to the presence of the native oxide layer around the nanostructures. In that case, when a liquid droplet, containing the analytes, is deposited onto the surface, it will entirely wet the surface. To better control the wetting properties and to avoid the droplet spreading, we have chemically functionalized the different surface with octadecyltrichlorosilane (OTS). For the N-SOI surfaces, we obtained WCA values of 125° after silanization. To compare, the WCA values obtained for the OTS-terminated from un-structured SOI and nanostructured crystalline Si (N-cSi) were 108° and 123°, respectively (Figure S6). Such WCA values of ~125° is known to be suitable for good and reproducible analyte deposition 25 .

III-3-2 Peptide MS detection

To go further, we have assessed the ability of our nanostructured SOI surfaces to perform LDI-MS detection of peptides. To do so, after their chemical modification by OTS, 0.5µL of peptide solution was deposited and let dried on each surface to be tested i.e. SOI, N-SOI and N-cSi. It has to be noticed that each deposition was done 3 times on same surface (3 spots) and 3 surfaces of each type were tested to check the reliability of the conducted experiments. Their performances were then evaluated in term of peptide discrimination (number of peptides detected/number of peptides in the mixture), signal intensities and laser fluence used. We started with the lower laser fluence and increase gradually until to obtain all peptides detected when possible. In figure S7 In figure S8 are presented mass spectra for peptide solution 1 deposited on N-cSi and N-SOI (BOx=0nm) surface and then detected by SALDI-MS. We can see huge difference between the 2 spectra with low signal intensities for N-cSi and only 2 on 4 peptides detected. N-SOI (BOx=0nm) showed higher signal intensity and good peptide discrimination (4 on 4 are detected). Despite the 2 surfaces were etched under the same conditions, they presented significant difference in terms of LDI-MS performances. One explanation could be that they have different thermal conductivities. Indeed, heat conduction in silicon is dominated by phonon transport. The N-SOI (BOx=0nm) thermal conductivity is reduced compared to that of N-cSi due to more important scattering mechanisms in the thin layer that are not present in the bulk material. Indeed, the reduction is more severe and extends to higher temperatures for thin layers than for bulk samples. Also important is phonon scattering on imperfections or grain boundaries, which exist in larger amount in SOI substrates compared to bulk material 34 . This phenomenon is amplified by structuring the thin layer thus enhancing the phonon scattering and lowering the thermal conductivity of N-SOI (BOx=0nm). A thermal conductivity for a porous silicon layer with a thickness of ~200 nm was measured to be 1.7

Wm -1 K -1 24 .

When buried oxide layer with various thicknesses (50, 100, 200 and 400 nm) are present between the nanostructured active layer and bulk Si, we can see that all N-SOI surfaces are able to detect all peptides present in the solution 1 (Figures S9 and6A). For each peptide from solution 1, almost same signal intensities' profile in function of BOx film thickness is observed with a significant increase of the signal intensities for N-SOI (BOx=200 nm) whereas a significant decrease is observed for BOx thickness of 400 nm (Figure S10).

From figure S10 and table S1, we can notice that for Glu 1 -Fibrinopeptide B and Neurotensin, their signal intensities and S/N are lower than for Des-Arg 1 -bradykinin and Angiotensin I. We know from previous study that SALDI-MS detection of peptides with masses superior to 1200 Da is more difficult to perform 25 . It has to be noticed that Glu 1 -Fibrinopeptide B is 5 times less concentrated in the mixture compared to others. Moreover, some peptides could be averse to be desorbed and ionized when present as part of mixtures even if they showed good MS data when they were analyzed as pure solution 26 . This could explain the high variation of signal intensities for Glu 1 -Fibrinopeptide B and Neurotensin. Nevertheless, this variation is lower for higher BOx thickness i.e. 200 and 400 nm.

We have also tested another peptide mixture (solution 2) composed by 5 peptides with higher masses (Figures 6B andS11). We can observe that only one type of N-SOI with BOx thickness of 200 nm is able to detect all the peptide with relevant intensities. Again, for all peptides considered within this work, the N-SOI (200nm) has shown the best LDI-MS performances in terms of peptide discrimination, signal intensities, S/N and spot to spot and surface to surface variations (Figure S12 and table S1). Concerning the LoD, the intensity profiles were recorded at 5pmol, 500 fmol and 50 fmol for each peptide solutions. Under our experimental conditions, an optimal detection at 10 -6 M including signal intensities and peptide discrimination for peptide solution 1 and 10 -5 M for peptide solution 2 were observed. However, at lower concentrations (such as 50 fmol), some peptides are very difficult to detect as mixture although they gave good data when analyzed in pure form.

In the contrary, we can notice that for N-SOI, with BOx thickness of 400 nm, a drastic decrease of signal intensities for all peptides considered in this study. A tentative explanation of this phenomenon could be drawn by the fact that by increasing the thickness of the thermal insulator, we decrease the thermal diffusivity (heat moves less rapidly). As a consequence, the surface will maintain a certain level of temperature (no rapid cooling) promoting the increase of peptide fragmentation thus decreasing the signal intensities of peptides (Figures S10 andS12). It was shown that there exists a tiny frontier between desorption and fragmentation and that the faster is the increase of the temperature, the lower is the fragmentation . Erreur ! Signet non défini. In the case of BOx=400 nm, the lower thermal diffusivity can keep the surface at a certain temperature thus promoting the peptide fragmentation.

As a control, a non-structured SOI surface (BOx = 200nm) was modified by OTS and then used for the LDI-MS detection of peptide solution 1 (Figure S13). The non-structured SOI surface only permitted the detection of 2 on 4 peptides at 40 % laser fluence with lower signal intensities values than all N-SOI independently of BOX thickness. More interestingly, even at low laser fluence, i.e. at 2% of laser fluence (instrument parameters), all peptides from peptide solution 1 have been detected when N-SOI (BOx=200 nm) surface was used (Figure S14). The use of low laser fluence will ensure the limitation of analytes fragmentation. In the contrary, when too high laser fluence was used, the thin nanostructured silicon active layer was completely ripped off as we can see on the SEM images (Figure S15).

We also performed the detection of model peptides spiked in serum to check the ability to detect peptides from complex solution. To do so, we made a solution 1 (1µL of Solution 1 in 24 µL of pure serum) and then deposit 0.5 µL of the mixture during 10 minutes. Then, the droplet was withdrawn and the remaining liquid let dried under ambient atmosphere prior MS detection. On figure 7, we can observe that all peptides from solution 1 were detected even when they are present in complex medium.

Finally, and to demonstrate that our N-SOI can also detect other type of compounds, antibiotics have been tested. It's well known that they are involved in current public health crisis because of their abusive use or misuse in food, animals or human treatments. In this extent, their SALDI-MS detection could be an interesting alternative of the classical current methods e.g. microbiological or immunoassays 35 . To do so, 4 antibiotics such as kanamycin m/z 484.5; novobiocin m/z 612.6; apramycin m/z 539.6; and erythromycin m/z 733.9 have been investigated at concentration of 100 µg/mL in pure or mixed solution (equimolar concentration) in presence of 10 mM NaCl. All antibiotics appeared in mass spectra in their protonated, sodiated or potassiated forms in both pure and mixed solutions (Figures S16 andS17).

IV-Conclusions

In this paper, we present the nanostructuration of the active top layer of SOI substrate via one-step MACE method. The fast preparation <1hour that includes 10 sec of etching and 30 min of chemical functionalization was sufficient to uniformly nanostructure the top active silicon layer of the SOI surface and use as SALDI surface. To investigate the influence of the BOx layer thickness on the LDI performances, we realize "home-made" SOI surface made of bulk silicon covered by various thicknesses of oxide and a thin layer of polysilicon (150 nm), both achieved via LPCVD.

The LDI-MS performances of the resulting surfaces were then assessed towards peptides and antibiotics mixtures and compared to references i.e. non-structured SOI surface and crystalline silicon structured using the same protocol. LDI-MS detection of peptides was successfully achieved on all tested surfaces with best results for N-SOI with a BOx thickness of 200 nm surface in terms of signal intensities, peptides discrimination and spot to spot and surface to surface variations.

We think that such nanostructured silicon layer is a very promising surface adaptable to any type of MALDI-ToF instrument and easy and fast to prepare. The good performances displayed at low laser fluence can also be useful in case of sensitive compounds limiting the risk of fragmentation. Also, application of such surface to mass spectrometry imaging (MSI) could be also considered. Indeed, the possibility to perform irradiation at sufficiently low laser fluence will help us preserving the biological content of the tissue by decreasing the ablation spot size, and so will increase the MSI resolution.

Other important parameters are the thermal conductivity and heat diffusivity that will control the area corresponding to the tissue vaporization. So, N-SOI surface could a very relevant alternative to stainless steel, ITO or other metal substrates used in MSI.

Figure captions

  ) Des-Arg1-Bradykinin m/z 904.05; Angiotensin I m/z 1296.51; Glu1-Fibrinopeptide B m/z 1570.61; Neurotensin m/z 1672.96 (solution 1) and ii) ACTH (clip 1-17) m/z 2094.46; ACTH (clip 18-39) m/z 2466.72, ACTH (clip 7-38) m/z 3660.19 and Insulin m/z 5734 (solution 2

  , it's shown the mass spectrum cleanliness obtained on N-SOI (Box 200 nm) surface compared to the one acquired with α-cyano-4-hydroxycinnamic acid (CHCA). N-SOI surfaces with different BOx thicknesses were first tested against 2 peptides solutions. Solution 1 composed by Des-Arg 1 -Bradykinin m/z 904; Angiotensin I m/z 1296; Glu 1 -Fibrinopeptide B m/z 1570; Neurotensin m/z 1673 (peptide mixture 1) and solution 2 composed by Angiotensin I m/z 1296; ACTH (clip 1-17) m/z 2094; ACTH (clip 18-39) m/z 2466; ACTH (clip 7-38) m/z 3660 and Insulin m/z 5734.

Figure 1 :

 1 Figure 1: Schematic representation of the whole procedure to nanostructure Si based surfaces, their chemical modification and their use as SALDI surfaces for analytes MS detection.

Figure 2 :

 2 Figure 2: SEM images comparing reference SOI and structured SOI achieved by MACE method for 10 sec of etching time.

Figure 3 :

 3 Figure 3: AFM characterization for the comparison of reference SOI and structured SOI surfaces achieved by MACE method for 10 sec of etching time.

Figure 4 :

 4 Figure 4: Reflectivity (R%) spectra measurements of the bare SOI and N-SOI (BOx=200 nm) (etching 10 sec) for wavelengths ranging from 250 to 800 nm.

Figure 5 :

 5 Figure 5: Photoluminescence spectra of the bare SOI and N-SOI (etching 10 sec) in the wavelength range of (400-1100 nm) at room temperature.

Figure 6 :

 6 Figure 6: MS spectra obtained from 2 peptides' solutions on N-SOI (BOx=200nm) substrate. Solution 1 (A) Des-Arg1-Bradykinin m/z 904; Angiotensin I m/z 1296; Neurotensin m/z 1673 at 50 fmol/µL and Glu1-Fibrinopeptide B m/z 1570 at 10 fmol/µL and solution 2 (B) Angiotensin I m/z 1296; ACTH (clip 1-17) m/z 2094; ACTH (clip 18-39) m/z 2466; ACTH (clip 7-38) m/z 3660 and Insulin m/z 5734 at 50 fmol/µL.

Figure 7 :

 7 Figure 7: Mass spectrum showing the detection of all peptides in human blood serum sample. Spike of solution 1 Des-Arg1-Bradykinin m/z 904; Angiotensin I m/z 1296; Neurotensin m/z 1673 at 50 fmol/µL and Glu1-Fibrinopeptide B m/z 1570 at 10 fmol/µL in blood serum.

Figure 1 :

 1 Figure 1: Schematic representation of the whole procedure to nanostructure Si based surfaces, their chemical modification and their use as SALDI surfaces for analytes MS detection.

Figure 2 :

 2 Figure 2: SEM images comparing reference SOI and structured SOI achieved by MACE method for 10 sec of etching time.

Figure 3 :

 3 Figure 3: AFM characterization for the comparison of reference SOI and structured SOI surfaces achieved by MACE method for 10 sec of etching time.

Figure 4 :

 4 Figure 4: Reflectivity (R%) spectra measurements of the bare SOI and N-SOI (BOx=200 nm) (etching 10 sec) for wavelengths ranging from 250 to 800 nm.

Figure 5 :

 5 Figure 5: Photoluminescence spectra of the bare SOI and N-SOI (etching 10 sec) in the wavelength range of (400-1100 nm) at room temperature.

Figure 6 :

 6 Figure 6: MS spectra obtained from 2 peptides' solutions on N-SOI (BOx=200nm) substrate. Solution 1 (A) Des-Arg1-Bradykinin m/z 904; Angiotensin I m/z 1296; Neurotensin m/z 1673 at 50 fmol/µL and Glu1-Fibrinopeptide B m/z 1570 at 10 fmol/µL and solution 2 (B) Angiotensin I m/z 1296; ACTH (clip 1-17) m/z 2094; ACTH (clip 18-39) m/z 2466; ACTH (clip 7-38) m/z 3660 and Insulin m/z 5734 at 50 fmol/µL.

Figure 7 :

 7 Figure 7: Mass spectrum showing the detection of all peptides in human blood serum sample. Spike of solution 1 Des-Arg1-Bradykinin m/z 904; Angiotensin I m/z 1296; Neurotensin m/z 1673 at 50 fmol/µL and Glu1-Fibrinopeptide B m/z 1570 at 10 fmol/µL in blood serum.

Acknowledgments

The Centre National de la Recherche Scientifique (CNRS), the Lille University, the Hauts-de-France and the CPER "Photonics for Society" are acknowledged for financial support. I. S. H. acknowledges the University Lille and DGA (Direction générale de l'armement, France) for PhD funding. The Renatech Network is also acknowledged for providing great support and technical conditions. This work was partly supported by the Ministry of Higher Education, Scientific Research and Technology of Tunisia.


Maxime Duris

Mohammed Guendouz

Nathalie Lorrain

Parastesh Pirasteh

Loïc Bodiou

Warda Raiah

Yannick Coffinier

Vincent Thomy

Joël Charrier
email: joel.charrier@univ-rennes1.fr

Vertical multilayer structures based on Porous Silicon layers for Mid-InfraRed applications

In this work, the fabrication of Porous Silicon Bragg reflector and vertical cavity on P + silicon substrate is investigated for applications in spectroscopic sensing in the Mid Infra-Red wavelength range. The complex refractive index of porous silicon layers is measured. Optical vertical devices are then fabricated and characterized by FTIR spectrophotometer. This work demonstrates the use of electrochemically prepared Bragg reflectors with reflectance higher than 95% and vertical cavity based on Porous Silicon layers operating in the Mid-IR spectral region (up to 8 µm). Experimental reflectance spectra of the vertical cavity structures are recorded as a function of air exposure duration after thermal annealing under nitrogen flux (N 2 ) and results demonstrate that these structures could be used for spectroscopic sensing applications in the Mid-IR field (2-8 µm) by grafting specific biomolecules on the porous silicon internal surface.

Introduction

Highly sensitive optical integrated sensors have received great interest during these last years [START_REF] Zinoviev | Silicon photonic biosensors for lab-on-a-chip applications[END_REF][START_REF] Soref | The past, present, and future of silicon photonics[END_REF][START_REF] Bogaerts | Silicon microring resonators[END_REF][START_REF] Fan | Sensitive optical biosensors for unlabeled targets: A[END_REF][START_REF] Densmore | Spiral-path high-sensitivity silicon photonic wire molecular sensor with temperature-independent response[END_REF][START_REF] Vos | Silicon-on-Insulator microring resonator for sensitive and label-free biosensing[END_REF][START_REF] Talebifard | Optimized sensitivity of Silicon-on-Insulator (SOI) strip waveguide resonator sensor[END_REF][START_REF] Armani | Label-free, single-molecule detection with optical microcavities[END_REF][START_REF] Grist | Silicon photonic micro-disk resonators for label-free biosensing[END_REF]. In particularly, the ability to rapidly detect, identify and monitor chemical or biological species is imperative for environmental monitoring, health monitoring and security applications. The detection of traces of (bio)-chemical molecules requires sample preparation procedures combined with sophisticated analytical tools which can detect, within an acceptable time, disease biomarkers, emerging pollutants, chemical warfare agents or toxic industrial chemicals with high sensitivity to really detect low concentrations with high selectivity not to be affected by other factors in the environment. The sensor should also have a fast response and recovery time whilst being portable, easy to operate and produce data that are easily interpreted. A range of optical, electrochemical, fluorescence, electrical and spectroscopic technologies have been developed for the detection of different kinds of chemical molecules [START_REF] Kim | A review of breath analysis for diagnosis of human health[END_REF]. Although useful information can be obtained and the technologies have progressed significantly, it is recognized that these methods are not adequate in terms of both spatial and temporal resolution. A major challenge is also the need to increase sensing sensitivity and detection reliability.

In recent years, Infra-Red (IR) spectroscopy has been proven to be a simple, reliable, relatively fast and cost-efficient technique for analyzing complex samples and determining their composition. The sensor produces a spectrum which represents the molecular absorption or transmission, creating unique molecular fingerprint of the sample. It is used in several research and industrial areas, from quality control to forensic analysis [START_REF] Mizaikoff | Waveguide-enhanced mid-infrared chem/bio sensors[END_REF][START_REF] Sieger | Toward On-Chip Mid-Infrared Sensors[END_REF][START_REF] Zheng | Development and field deployment of a mid-infrared methane sensor without pressure control using interband cascade laser absorption spectroscopy[END_REF][START_REF] Li | Development and measurement of a near-infrared CH4 detection system using 1.654 μm wavelength-modulated diode laser and open reflective gas sensing probe[END_REF][START_REF] Song | Interband cascade laser based midinfrared methane sensor system using a novel electrical-domainself-adaptive direct laser absorption spectroscopy (SA-DLAS)[END_REF][START_REF] Dong | Doublerange near-infrared acetylene detection using a dual spot-ring Herriott cell(DSR-HC)[END_REF][START_REF] Gattass | Infrared glass-based negative-curvature antiresonant fibers fabricated through extrusion[END_REF][START_REF] Rodrigo | Mid-Infrared Plasmonic Biosensing with Graphene[END_REF][START_REF] Singh | Mid-IR absorption sensing of heavy water using a silicon-on-sapphire waveguide[END_REF][START_REF] Lochbaum | On-Chip Narrowband Thermal Emitter for Mid-IR Optical Gas Sensing[END_REF][START_REF] Kumari | Silicon-on-nitride slot waveguide: A promising platform as mid-IR trace gas sensor[END_REF]. The IR region of the electromagnetic spectrum can be divided into the near, mid and far IR regions. The most common wavelength range for detection applications lies in the Mid Infrared (Mid-IR) region and particularly ranges between 2.5 μm and 15 μm (4000 -670 cm -1 ). While current spectroscopic systems are based on bulk optical components, the resulting systems are large, expensive and not rugged. Photonic integrated circuits provide a route to small, low-cost and very rugged optical systems and could therefore be a game-changer for IR spectroscopic sensor systems. However, most integration technologies developed up to now are geared toward sub 2µm wavelength operation, which does not include the wavelength range of most interest for spectroscopic sensing, namely the mid-infrared.

The ability to tailor Porous Silicon (PSi) refractive index and layer thickness by controlling porosity and anodization etching time [START_REF] Canham | Handbook of porous silicon[END_REF][START_REF] Canham | Handbook of porous silicon[END_REF][START_REF] Arshavsky-Graham | Analytical Chemistry Porous Silicon-Based Photonic Biosensors[END_REF] makes it especially attractive for optical applications. Optical multilayer devices such as Bragg reflectors or micro-cavities have already been implemented from PSi layers notably for Visible and Near Infra-Red (NIR) wavelength ranges [START_REF] Wang | Engineering the Properties of Polymer Photonic Crystals with Mesoporous Silicon Templates[END_REF][START_REF] Martin | Matrix metalloproteinase sensing via porous silicon microcavity devices functionalized with human antibodies[END_REF][START_REF] Chan | Porous silicon microcavities for biosensing applications[END_REF][START_REF] Chan | Identification of gram negative bacteria using nanoscale silicon microcavities[END_REF][START_REF] Ouyang | Macroporous silicon microcavity for macromolecule detection[END_REF]. Two constant current densities are used for the High Porosity (HP) and Low Porosity (LP) layers with low and high refractive indices (n HP and n LP ) respectively for the fabrication of these optical devices. The etching times are adjusted to produce an optimized Bragg reflector or cavity with the resonant wavelength in the chosen spectral range. Moreover, PSi is also biocompatible material and the functionalization of its internal surface by grafting molecules to detect is interesting for surface detection due to its high specific surface that is very important and up to 800 m 2 /cm 3 [START_REF] Kumari | Silicon-on-nitride slot waveguide: A promising platform as mid-IR trace gas sensor[END_REF][START_REF] Canham | Handbook of porous silicon[END_REF]. The implementation of a Mid-IR silicon photonics transducer with broad Mid-IR transparency (up to 8 µm by taking into account Si transparency) is a challenge that could find applications in spectroscopic sensing and environmental monitoring. This paper demonstrates the fabrication of vertical Porous Silicon multilayer structures on Si substrates and their potential in the Mid-IR wavelength range notably near the cut-off band of Si due to its absorption at 8 µm. That is why the fabrication of Bragg reflector and vertical cavity on P + silicon substrate is investigated for applications in spectroscopic sensing in the (Mid-IR) wavelength range

Experimental conditions

The PSi layers were obtained by electrochemical anodization of p-type doped (100)oriented silicon wafer (4-6 m.cm) at room temperature. The electrolyte was composed of HF(50%):ethanol: deionized water (2:2:1). The refractive indices and porosities of PSi single layers were studied by reflectometry in the near-and mid-infrared spectral ranges. Scanning electron miscroscopy (SEM) was used to obtain the thickness of PSi layers. The thickness of each porous layer was controlled by the anodization etching time. Current densities of 20 and 90 mA/cm 2 were used for the low porosity (LP) and high porosity (HP) respectively. . These current densities were chosen to obtain a high index contrast. The infrared reflectance spectra were measured by a Perkin Elmer Spectrum 100 FTIR spectrophotometer from 1.5 µm to 15 µm with a resolution of 4 cm -1 . The calibration is performed by a measurement using a gold mirror with a calibrated reflection equal to 97.5% between 1.5 and 25 µm. The calibration protocol of the spectrometer takes into account absorption from the environment in particular absorption fringe of CO 2 , water vapor and methane. From reflectance spectra, the complex refractive index of the fabricated single layers was deduced by fitting spectral fringes using the effective medium theory (Bruggeman model) [START_REF] Theiss | Optical properties of porous silicon[END_REF][START_REF] Aspnes | Optical properties of thin films[END_REF]. Some samples were annealed under N 2 for 1 hour in order to desorb the specific surface of the porous layers and to remove the contaminants.

From the refractive index and the anodization etching rate knowledge of the two studied PSi layers, the design of vertical optical devices was performed using the transfer matrix method in order to determine the best combination of parameters for Bragg reflectors and cavities (reflection bandwidth, resonance wavelength and maximum reflectance).

Results and discussion

The porosities and associated Mid-IR refractive indices have been determined from reflectivity measurements of two single PSi layers whose fabrication parameters, given above in the experimental part, correspond to those of the multilayer structures. From the thicknesses measured by scanning electron microscopy, porosities and then complex refractive indices were determined by the adjustment of the calculated reflectance spectra of each PSi layer with experimental ones using Bruggemann model [START_REF] Charrier | Bragg reflector formed on oxidized porous silicon[END_REF][START_REF] Najar | Ultra-Low Reflection Porous Silicon Nanowires for Solar Cell Applications[END_REF][START_REF] Lorrain | Functionalization control of porous silicon optical structures using reflectance spectra modeling for bio sensing applications[END_REF]. The experimental reflectance spectra are shown in Figure 1 (a) for the two PSi single layers obtained with the current densities of 20 and 90 mA/cm 2 . The experimental fringes are well visible until 8 µm before the silicon absorption in the Mid-IR range. The adjustment of the experimental spectra was carried out for wavelengths ranging from 2 to 8 µm. A good fit between experimental and theoretical spectra was obtained as illustrated in Figure 1 (b). The complex refractive indices and the extinction coefficients for the two PSi porosities are reported in Figures 2 (a It should be noted that the extinction coefficient of porous layers is low in 1550 -8000 nm spectral region as expected. The real refractive index contrast between the HP and LP films was found to be 0.88 at 1550 nm and 0.62 at 7 µm. The porosity values are respectively for the low and high porosities equal to 33 and 61 % and the extinction coefficient k is higher for the layer of lower porosity as indicated in Table 1 due to the contribution, of the Si absorption that is more important for this low porosity value. In order to avoid the need for thick films in the multilayers devoted to mid-IR applications, a large index contrast was targeted between the layers forming the pattern. These values are therefore very promising for targeted applications considering the low absorption of the PSi layers. structure of the porous layers is well observable in Figures 3(a) and(d). This morphology is characteristic of mesoporous layers that are fabricated from P + Si substrate [START_REF] Kumari | Silicon-on-nitride slot waveguide: A promising platform as mid-IR trace gas sensor[END_REF][START_REF] Canham | Handbook of porous silicon[END_REF]. The interfaces between individual layers as well as the surfaces of the multilayered structures were found to be smooth without any defect during the fabrication of thin film stacks (Figures 3(b), (c) and(d)). The achieved reflectivity is determined by the number of layer pairs N and by the refractive index contrast n between the layers. The Bragg reflector is characterized by its central wavelength  0 (at normal incidence) and by the reflection bandwidth  which is determined mainly by the index contrast. These two parameters are defined respectively by the relationships (1) and [START_REF] Soref | The past, present, and future of silicon photonics[END_REF].

  LH LH LP LP t n t n   2 0  (1) n n 2 0       (2) with 2 n n n n n n HP LP HP LP     
Where n HP (LP) and t HP (LP) are the refractive index and thickness of high (low) porosity layers.

The thicknesses of pattern were fixed to obtain Bragg reflector and cavity structures centered at 6.5 µm. The experimental responses of the Bragg reflector is reported in Figure 4(a). The spectrum is centered close to 6560 nm. The experimental wavelength bandwidth is about 120 nm and the maximal reflectance is around 97%. The experimental spectrum is in close agreement with the theoretical spectrum taking the absolute errors into account. The absolute errors of measured refractive index and thickness of layers are 0.005 and 50 nm respectively. This work demonstrates that it is possible to fabricate a Bragg reflector with good optical qualities in the Mid-IR range from P + silicon substrates.

With the same pattern, a vertical cavity was fabricated with a resonance wavelength equal to 6550 nm. The experimental reflectance spectrum of the cavity is illustrated in figure 4(b). The resonance peak is close to the target value. Moreover, the reflectance value of the device of the bandwidth is as high as 97%.

Another merit factor is the quality factor of the structure. In order to study this parameter, we can use the Fabry-Perot classical model [START_REF] Charrier | Merit factor theoretical study of optical biosensors based on porous silicon[END_REF]. The quality factor is given by the following relationship:

    L R L R L n Q             exp 1 2 exp 2 0 0 0 (3)
where n0 is the refractive index of the defect layer that separates the two Bragg mirrors (n0= 2.07), L its thickness 1.75 µm), 0 the expected resonant wavelength equal to 6560 nm,  its absorption coefficient (2.3 cm -1 ) [START_REF] Pirasteh | Optical loss studies of porous silicon and oxidized porous silicon planar waveguides[END_REF] and R the reflectance value equal to 97%by assuming that the two Bragg mirrors are similar.

For the elaborated microcavity, the theoretical quality factor is equal to 112 and from the experimental spectrum, the quality factor was estimated at around 60. However, we noted a degradation of the quality factor as a function of time when the sample was exposed to air as shown in figures 5 . After thermal annealing under N 2 at 300 °C for one hour, the spectrum of initial experimental reflectance has been restored due to the desorption of the specific surface of the sample. Several experimental reflectance spectra were recorded as a function of exposure duration to air after thermal annealing at 300°C under N 2 . These spectra are reported in Figure 5 (a). We note a degradation of the cavity quality factor that is all the greater the longer the exposure of the micro-cavity to the ambient air. After a new heat treatment under N 2 of the cavity exposed to air for 30 minutes, the quality factor is again improved. This quality factor reduction is better illustrated in Figure 5 (b) by showing the spectra near the resonance wavelength to emphasize the resonance FWHM and reflectance value increase. The degradation of the quality factor is due to the adsorption on PSi internal surface of molecules contained in the ambient air. Indeed, molecules such as water vapor or methane have absorption lines in the studied wavelength range, around of 6550 nm.

A red shift is noted for the resonance peak and the width of the peak increases strongly due to the molecule adsorption with air exposure time. The initial reflectance is almost obtained after a new annealing under N 2 of the 30 minutesair exposed cavity (green curve in figure 5 (b)). The difference between the two spectra (just annealing and new annealing after 30 minutes exposure) is due to the difference time between the FTIR measurement and the sample exit from the oven. Reflectance spectrum overlaps absorption bands of air molecules and their detection is enhanced by the cavity [START_REF] Martin | Matrix metalloproteinase sensing via porous silicon microcavity devices functionalized with human antibodies[END_REF][START_REF] Chan | Porous silicon microcavities for biosensing applications[END_REF][START_REF] Chan | Identification of gram negative bacteria using nanoscale silicon microcavities[END_REF][START_REF] Ouyang | Macroporous silicon microcavity for macromolecule detection[END_REF]. So by a specific functionalization of porous silicon [START_REF] Arshavsky-Graham | Analytical Chemistry Porous Silicon-Based Photonic Biosensors[END_REF], these vertical optical structures could be used as transducers to detect low concentrations of target molecules by choosing the good resonance wavelength, in the Mid-IR range, and relative to its absorption band. However, such microcavities should be encapsulated in a neutral atmosphere so that a microfluidic system only interacts with a fluid containing molecules to be detected.

Conclusions

In conclusion, vertical multilayer stackings based on PSi layers were fabricated for Mid-IR wavelength range. From the experimental reflectance spectra of PSi thick layers, real and imaginary parts of the refractive index of these films of different porosities were obtained in the Mid-IR wavelength range. Bragg mirror and vertical cavity were then fabricated for notably near the cut-off band of Si due to its absorption at 8 µm. Multilayer structures based on PSi layers made from P + Si substrates were then implemented. Maximal reflectance around 95% and a quality factor equal to 50 were obtained respectively for Bragg mirror and vertical microcavity. These optical devices could be used as highly sensitive transducers thanks to the PSi layers opened pores and wide internal surfaces. The first results obtained from these PSi multilayer structures can now serve as functionalized transducers by grafting biomolecules to develop spectroscopic sensing applications in the Mid-IR field (2-8 µm).

Funding and acknowledgments

Fig. 1 .

 1 Fig. 1. (a) Experimental Mid-IR reflectance spectra of the PSi layer of high and low porosities and of the Si substrate. (b) Experimental and fitted reflectance spectra of the high porosity PSi layer

Fig. 2 .

 2 Fig. 2. Variation of (a) PSi refractive index and (b) PSi extinction coefficient as a function of the wavelength for the high and low porosity layers.

Fig. 3 .

 3 Fig. 3. SEM micrographs (cross sections) of (a) monolayer of PSi, (b) Bragg reflector structure and (c,d) vertical microcavity at some different scales. Some examples of SEM micrographs of single layers (Figure 3(a)) and prepared multilayered structures such as Bragg mirror (Figure 3(b)) or vertical cavity (Figures 3(c) and figure 3(d)) are presented in Figures 3. The columnar structure of the porous layers is well observable in Figures3(a) and (d). This morphology is characteristic of mesoporous layers that are fabricated from P + Si substrate[START_REF] Kumari | Silicon-on-nitride slot waveguide: A promising platform as mid-IR trace gas sensor[END_REF][START_REF] Canham | Handbook of porous silicon[END_REF]. The interfaces between individual layers as well as the surfaces of the multilayered structures were found to be smooth without any defect during the fabrication of thin film stacks (Figures3(b), (c) and (d)).

  ) and figure 3(d)) are presented in Figures 3. The

  columnar

Fig. 4 .

 4 Fig. 4. Experimental reflectance spectra of (a) Bragg reflector and (b) vertical caviy structures.

Fig. 5 .

 5 Fig. 5. (a) Evolution of the experimental reflectance spectra of the same microcavity structure as a function of exposure duration to air after annealing at 300°C of the structure under N2. (b) Evolution of the resonance wavelength of the same microcavity structure as a function of exposure duration to air after annealing of the structure under N2.

Table 1 . Physical parameters of the lower and higher PSi layers

 1

	Refractive	index	and
	extinction coef.		@ 1550 nm	@ 7 µm
	Name of	Value	of
	layer	porosity (%)	n	k	n	k
	LP	33 ± 6		2.35 ± 0.05	0.035 ± 0.004	2.07 ± 0.05	0.027 ± 0.004
	HP	61 ± 4		1.47 ± 0.03	0.01 ± 0.002	1.45 ± 0.03	0.008 ± 0.002

Acknowledgments

Scanning Electron Microscopy imaging and porous silicon fabrication were performed in the CCLO -Renatech clean room facilities of Institut Foton.

Funding

This work was supported by the French national research agency under MID-VOC ANR project (ANR-17-CE09-0028-01). Equipments funding of Institut Foton were partly provided by the CPER Sophie.