Carmen Ciotonea
email: carmen.ciotonea@univ-lille.fr

Alexandru Chirieac

Brandusa Dragoi

Jérémy Dhainaut

Maya Marinova

Stephane Pronier

Sandrine Arii-Clacens

Jean-Philippe Dacquin

Emil Dumitriu

Adrian Ungureanu
email: aungureanu@tuiasi.ro

Jeremy Dhainaut

Sébastien Royer

Playing on 3D spatial distribution of Cu-Co (oxide) nanoparticles in inorganic mesoporous sieves: Impact on catalytic performance toward the cinnamaldehyde hydrogenation

Keywords: transition metal oxide, mesoporous support, supported catalyst, bimetallic catalyst, hydrogenation reaction

• INTRODUCTION

The design of supported metal nanoparticle catalysts has attracted considerable interest in various catalytic applications of heterogeneous catalysis [1] where the necessity of using more than one type of metallic active sites is often requested due to the complexity of the studied reactions [2]. Therefore, numerous research groups focused their work on the preparation methods and/or applications of (bi)metallic supported nanoparticles (NPs) as catalysts in various (de)hydrogenation processes involved in important fields such as fine chemistry, environmental pollution control, energy and so on [3]. As reported in literature, the insertion of a second active metal to the catalytic phase has the capacity to boost the properties of the first metal due to the modification of its electronic properties [1][2][3][4]. In this line, non-noble (bi)metallic supported nanoparticles (NPs) (e.g.Ni, Cu, Co…) offer a variety of technical and commercial advantages, which make them feasible substitutes to the more expensive supported noble metal NPs (e.g. Pt, Pd, Ir, Au) [4]. The particular NPs morphology (shape and size) and their chemical composition (affecting the metal-support and metal-reactant interactions) are the main variables which affect the catalytic activity of supported bimetallic NPs. Consequently, by optimizing the dispersion and accessibility of the active metal species, greatly enhanced catalytic performances of supported NPs (i.e. activity and selectivity) have been reported with respect to their bulk counterparts. Those observations can be explained by larger surface-to-volume ratios of NPs (i.e. higher dispersion) associated with an increased number of defect sites, kinks, steps, edges and corners [5]. Therefore, considerable research efforts are focused nowadays on the development of new preparation methods which can produce such advanced catalytic nanomaterials via both physical and chemical approaches [2][3][4][5]. To reach the production of NPs with similar physico-chemical properties, it is desirable to use a support with ordered array of pores and high surface area, in order to get a better control in the NPs deposition and an easier understanding of the stabilisation processes that took place during the preparation and activation of the catalyst. Among the candidates fitting all the above conditions, the SBA-15 presents a highly ordered mesoporous siliceous skeleton with advanced characteristics comparable to other silica homologues [6],

and makes it ideal to be used as host structure: (i) high specific surface areas (600-1000 m 2 g -1 ); (ii) high (hydro)thermal stability with respect to its MCM-41 counterpart, well defined pore structure formed of hexagonal array of large and tunable cylindrical pores (diameter 5-10 nm) which allow the confinement of NPs and also a very good thermal stabilisation towards particle sintering [7]. Also a secondary porosity can be produced and, in this case, small particles may be immobilized onto the support porosity, resulting in a material with enhanced properties in term of stabilisation and dispersion. [8][9][10][11] By far the most frequently used method in both academy and industry for the preparation of supported NPs is the incipient wetness impregnation technique (IWI), due to some essential sustainability-related characteristics such as versatility [12][13][14], convenience, low amount of required solvent, availability and low cost of metal sources, high purity of metal precursor phases and limited production of waste. Being easy to obtain, the principle of IWI method [28] is to mix an aqueous metal source with a porous support, followed by drying and thermal decomposition and/or reduction of the precursors to obtain the oxide phase [15].

Bi-metallic materials based on Cu and Co NPs are very interesting to be studied as a potential synergic catalytic system. The copper is located in periodic table, within the 11th transition metal group. The specificity of this group is that the "d band" is fully occupied. Thus, copper is expecting to manifest less affinity for accepting electrons and consequently, we should observe poor selectivity under the hydrogenation reaction. Interestingly, copper, as well as the other elements from the 11th group are showing very good selectivity toward the hydrogenation of C=O band of α,β-unsaturated aldehydes [4]. On the other side, cobalt is located in the 9 th group, thus partially filled d band, and is showing excellent selectivity toward the production of unsaturated alcohol from the hydrogenation of citral, acetophenone and cinnamaldehyde [16].

Regarding the literature, Cu and Co bi-metallic materials were deposited onto various supports such: SiO2 [17][18], Al2O3 [19] TiO2 [20][21] porous carbon [22], etc. In an attempt to summarize the above observations, three situations can be found on solids based on Cu-Co supported on silica: formation of single monometallic oxides of i) large CuO NPs; ii) the spinel Co 2+ 1Co 3+ 2O4 NPS or iii) combination of both Co and Cu within the bimetallic spinel, with general formula CuxCo(3-x)O4 [23][24][25]. The Co3O4 spinel belonging to the space group Fd3m has oneeighth of the tetrahedral sites occupied by Co 2+ cations, while one half of the octahedral sites are occupied by Co 3+ cations, consequently, the insertion of Cu 2+ within the spinel will take place in the tetrahedral position of Co 2+ as follow CuxCo(1-x)Co2O4, where x can have only values inferior to 1. [23,[25][26][27][28].

Albeit IWI is often considered as a straightforward way to produce metal NPs, numerous results are showing a heterogeneous distribution of the single oxides and high amount of external large particles, thus there is space to improve the IWI procedure [29].

Accordingly, our undertaking research is motivated by the possibility to optimize the procedure of the IWI method in order to outcome a general, controlled/reproducible and facile impregnation approach to prepare transition (bi)metallic NPs supported on SBA-15. Herein, we report the synthesis of CuCo NPs stabilized within the ordered mesoporosity of SBA-15 silica through a modified incipient impregnation method including mild drying step (25°C, ±2°C). In this context, our objectives are the exploitation of factors that affect the physico-chemical properties of non-noble bimetallic nanoparticles (Cu and Co) supported on SBA-15 for enlarging the geometric and electronic features such as the particle size and morphology, the nature of the two metals and their electronic behavior influencing the metal-reactant interactions. For the latter properties, the reactivity of CuCo NPs supported onto SBA-15 has been evaluated towards the hydrogenation of cinnamaldehyde (CNA). The chemoselective hydrogenation of α,β-unsaturated aldehydes is very sensitive to the behaviour of the divided phases [4]. Within this family of reactions, the selective production of cinnamyl alcohol (CNOL) by hydrogenation of cinnamaldehyde (CNA), is highly sought for, because the unsaturated alcohol is an important additive in pharmaceuticals, agrochemicals, flavours and fragrances. For that reason, hydrogenation of CNA draws increased attention in the fine chemical sector as one of the current challenges to be tackled from both fundamental and industrial viewpoints. In industry the production of cinnamyl alcohol is realised by Meerwein-Ponndorf -Verley reduction, while the catalytic processed is realised by using OsC as catalyst [30][31]. The main difficulty arises from the hydrogenation of C=C bond that is thermodynamically and kinetically favoured [32] and thus the saturated hydrocinnamaldehyde (HCNA) is usually the major by-product. Furthermore, a competition reaction took place for the hydrogenation to the second unsaturated bond from either the unsaturated alcohol or the saturated aldehyde already formed, subsequently the saturated hydrocinnamyl alcohol (HCNOL) is produced (Scheme 1, SI file).

Until now, Cu-Co supported on silica were previously used in the hydrogenation of α,β-unsaturated aldehydes, with good conversion up to 50 % and high selectivity. The necessity of the copper into the bimetallic system is due to the fact that the latter is favouring the hydrogen dissociation, while the adsorption of the cinnamaldehyde will take place onto the active sites of M 2+ (with M 2+ = Cr, Co, etc) via η1 on top adsorption, thus the presence of the sites formed by Cu 0 M 2+ are preferred to be used for the CNA hydrogenation toward the enhancement of unsaturated alcohol [4,33].

In our method, the main results dealing with the chemoselectivity of the reaction will be discussed and linked with geometric and electronic features which are in close relationship with the physico-chemical properties of the porous host in which the bimetallic NPs are stabilized.

• EXPERIMENTAL

Catalyst preparation

Materials: All chemicals required for the synthesis of the mesoporous silica support and derived supported metal catalysts were used as purchased: tetraethylorthosilicate (Si(OC2H5)4, TEOS-98 wt.%, Sigma-Aldrich), non-ionic triblock copolymer Pluronic P123 (poly(ethyleneoxide)-block-poly-(propyleneoxide)-block-poly(ethyleneoxide)block), EO20PO70EO20, molecular weight of 5800, BASF Corp.), hydrochloric acid (HCl, 37 wt.%, Sigma-Aldrich), cobalt nitrate (Co(NO3)2•6H2O, 98 wt.%, Sigma-Aldrich) and copper nitrate (Cu(NO3)2•3H2O, 98 wt.%, Sigma-Aldrich). For the catalytic runs, the chemicals were also used as received: trans-cinnamaldehyde (C6H5CH▪CHCHO, 98 wt.%, Merck) as reagent, propylene carbonate (C4H6O3, 99 wt.%, Sigma-Aldrich) and isopropanol (C3H8O, Merck 99 wt.%) as solvents.

Support synthesis: Ordered mesoporous SBA-15 support was prepared by using chemical acidic conditions [34] of HCl (1.6 M) by using Pluronic P123 (4 g). The resulting mixture was treated at 40 °C. After complete dissolution of the structure directing agent, the silica source, TEOS (8.5 g), was added dropwise to the solution, followed by magnetic stirring for 24 h at 40 °C. A hydrothermal treatment for 48 h at 100 °C was used for consolidation of the gel. The hybrid solid was recovered by filtration, washed, and dried under air. In order to obtain free porosity, the material was calcined at 550 °C for 6 h (in a muffle furnace, using a heating ramp of 1.5 °C min -1 ).

IWI method:

Transition metal oxide-loaded SBA-15 samples were prepared by incipient wetness impregnation under mild drying conditions [15]. The volume of TM precursor aqueous solution added for IWI-MD is equal to Vp measured by N2 physisorption for the parent SBA-15 used. The equivalent of aqueous metal nitrate precursors solutions were mixed for achieving a constant loading of zero-valent metal components of 5 wt.% and variable Cu/M weight ratio (in the range of 0 -5; M = Cu + Co). The solution containing solids was gently dried under air at 25 ± 1 °C for 48 h. As previously showed by our group, dry step performed at room temperature is improving the dispersion of the NPs compared to classical IWI (drying step at 100°C). The powders were thereafter calcined under air at 500 °C for 6 h (heating ramp of 1.5 °C. min -1 ) to obtain the oxidic form of catalysts. To avoid the "container effect" during calcination, open crucibles and similar sample depths were used. After calcination, the solids were stored under ambient conditions in sealed containers. The monocomponent and bicomponent samples were labelled according to their chemical composition as CuCox:y, where x:y are namely Cu and Co in relative proportions (x:y = 5:0, 1:4, 1:1, 4:1, 0:5).

Materials characterisation:

Materials were systematically characterized in their oxidized and/or reduced state using XRD at low and ex-situ and in-situ large angles, nitrogen physisorption, TEM-EDX, TPR, XPS and ICP-OES. Catalytic tests were performed after a reduction treatment at 500 °C in order to evaluate the beneficial impact of the metallic preparation on catalytic potential of the final materials.

ICP-OES analyses were performed using a sequential scanning inductively coupled plasma with an optical emission spectrometer (Perkin Elmer Optima 2000 DV) to determine the chemical composition of the samples. Before analysis, a known amount of sample was dissolved in a diluted HF-HCl solution under microwave heating.

Results showed that transition metal final loading is slightly higher than the expected 5 wt.% value.

N2-physisorption isotherms were recorded at -196 °C on an Autosorb 1-MP automated gas sorption system (Quantachrome Instruments). Before analysis, the samples were outgassed under dynamic vacuum at 350 °C for 3h. Textural properties were determined from the adsorption/desorption isotherms by using the Autosorb 1 software. BET surface area was determined using the multipoint BET algorithm in the P/P0 range of 0.1-0.25. The t-plot method was applied to quantitatively determine the micropores volume and surface areas (de Boer statistical thickness P/P0 of 3.8-6.5 Å). The mesopore size distribution was determined by NLDFT equilibrium algorithm applied for cylindrical pores.

Micrographs (transmission electron microscopy, TEM) on oxidized samples were recorded on a JEOL 2100 UHR instrument (operated at 200 kV with a LaB6 source and equipped with a Gatan Ultra scan camera), with a resolution of 0.19 nm, and equipped with an Energy Dispersive X-ray Spectroscopy (EDXS) detector. All the samples were embedded in a polymeric Epoxy resin (spurr) and cut into sections of 50 nm using an ultramicrotome equipped with a diamond knife. Cuts were deposited on gold grids in order to get proper information about the copper content from the samples when analysing. The images on samples reduced at 750 °C were collected on a TITAN Themis 300 S/TEM with a probe aberration corrector and monochromator, allowing spatial resolution of 70 pm and energy resolution of 150 meV. The microscope was equipped with a super-X windowless 4 quadrant SDD (silicon drift detector) detection system for the STEM-EDX mapping and with several annual dark field detectors.

The experiment was performed with a 0.5 nm probe size, a convergence angle of 21 mrad and a probe current of approximately 100 pA. For the high angle annular dark field (HAADF), images collection angles have been between 50 and 200 mrad. X-ray diffraction was performed using PANalytical Empyrean X-ray diffractometer, in Bragg Benttano configuration, with Cu K radiation (λ = 1.54184 Å). For small-angle analysis, diffractograms recorded for the data were collected in the 2θ ranging from 0.5° to 5°, with a step size of 0.01°. Wide-angle analyses were performed for a 2θ range, laying between 10° and 80° with a step size of 0.05°. Mean crystal size was calculated using the Scherrer equation: Dhkl = K.λ/b.cosθ, where K is a structure constant (0.9, for spherical crystals); λ is the incident ray wavelength; b is the peak width at half height after correction for instrumental broadening; and θ is the Bragg angle. Powder XRD: patterns were recorded on a Bruker D8 ADVANCE X-ray diffractometer equipped with a VANTEC-1 detector, using a Cu Kα radiation (λ =1.54184 Å) as X-ray source. The calcined samples were placed on a kanthal filament(FeCrAl) cavity, and were subjected to thermo-programmed reduction under a flow of 3.0 vol.% H2 in He (30 mL min-1) from 30 °C to 500 °C (heating ramp of 5 °C min -1 ). The in situ diffractograms were recorded at definite temperatures in the 2θ range from 15 to 70° with a step of 0.05°. Crystal phase identification was made by comparison with ICDD database.

Temperature-programmed reduction under H2 (H2-TPR) was conducted on an Autochem chemisorption analyser (Micromeritics), equipped with a TCD. The consumption of H2 was driven from the TCD signal after calibration. Prior to the H2-TPR run, the solid was activated up to its calcination temperature (from room temperature to 500 °C, heating ramp of 5 °C min -1 , and isothermal time of 1 h) under simulated air at a total flow rate of 30 mL min -1 . After cooling down to 50 °C, the H2 containing flow was stabilized (30 mL min -1 , 3.0 vol.% H2 in Ar) and the temperature-programmed reduction was performed (from 50 °C to 900 °C, with a temperature ramp of 5°C min -1 ).

X-ray photoelectron spectroscopy (XPS) analyses were performed with a Kratos Axis Ultra, equipped with a dual Mg/Al anode and a monochromatized Al Kα source (1486.6 eV). The samples were analyzed as pellets, mounted on a double-sided adhesive tape. The pressure in the analysis chamber was in the range of 10-9 mbar during data registration. The binding energies were referred to the C1s component positioned at 284.8 eV. Atomic concentrations were calculated from peak intensity using the sensitivity factors provided with the Casa XPS software (version 2.3.16 PR 1.6). The binding energy values are quoted with a precision of ±0.15 eV and the atomic percentage with a precision of ±10%.

Hydrogenation reaction of cinnamaldehyde

Performance in catalysisat atmospheric pressure: liquid phase hydrogenation of cinnamaldehyde was performed in a three phase system, in a thermostatic three-neck glass reactor equipped with reflux condenser, hydrogen bubbler and magnetic stirrer. The following reaction conditions were applied: temperature of 150 °C, 1.05 g of trans-cinnamaldehyde, 25 mL of propylene carbonate, 0.265 g of activate catalyst, hydrogen flow of 1 L h -1 , and stirring rate of 900 rpm.

-at high-pressure: liquid phase hydrogenation of cinnamaldehyde reaction was performed by using a Parr reactor under the following typical conditions: 1 mL cinnamaldehyde, 40 mL isopropanol as a solvent, 0.133 g catalyst, 10 bar H2 and reaction temperature of 130 °C.

Prior to the catalytic runs, the calcined samples were gently crushed and sieved to obtain a granulometric fraction lower than 0.126 mm. Transition metal oxide phases, obtained after calcination of the materials, were then reduced to the metallic form at 350 °C and 500 °C for 10 h (heating ramp of 6 °C min -1 ) under H2 flow (1 L. h -1 ). Preliminary tests made with different granulometric fractions and loadings of catalyst, and different stirring rates, disclosed no diffusional limitations under the selected reaction conditions. Aliquots of reaction mixture were withdrawn periodically and analysed by gas chromatography (HP 5890 instrument, equipped with a DB-5 capillary column and a FID detector). The identification of the reaction products was achieved from the known retention times of pure compounds and occasionally by GC-MS (Agilent 6890N system equipped with an Agilent 5973 MSD detector and a DB-5-ms column). The conversion of cinnamaldehyde and selectivities in the different hydrogenation products were calculated by taking into account the FID response factors for each compound.

RESULTS AND DISCUSSION

Surface area and porosity

The adsorption/desorption isotherms and related pore size diameter distributions of supported CuCo NPs on SBA-15 were registered by N2 physisorption at -196 °C and are presented in Figure 1. Corresponding textural properties are summarized in Table S1. For all freshly calcined materials, isotherms of type IV with H1-type hysteresis loops are easily observed, which evidence the presence of highly ordered mesoporous silica type SBA-15 with cylindrical channels. Parallel branches of the hysteresis loop are indicating a uniform periodic arrangement of the mesopores, as it can be also noticed by the derived pore size distribution (figure 1). A latency can be observed in closing of the hysteresis which arises from the presence of NPs inside the channels. Such pore blockages induce the formation of ink-bottle shape pores, visible onto the hysteresis (i.e. on the desorption branch side) at pressures ranging from 0.58 to 0.78 [35]. The pore size distribution, from NL-DFT model, is showing an uniform pore size distribution of 8.4 nm and are in close agreement with the low angle XRD patterns (Figure S1). The presence of NPs inside the mesoporosity can be first indicated by the formation of smaller mesopores of 7.3 nm. Further evidence of the NPs confinement is highlighted by the decrease of the specific surface area of the CuCo materials (from 760 m².g -1 to 575 m 2 g -1 ) with respect to the parent SBA-15 (880 m².g -1 ).

Structure and Morphology

For all samples, low angle XRD analysis is revealing typical peaks corresponding to the planes of 100, 110 and 220 (Figure S1). All planes are representative of a high periodic arrangement of the channels of the silica type SBA-15 and is associated to p6mm symmetry. The presence of the NPs inside the mesoporosity is as well confirmed by the shift at higher 2θ for d100 peak which appears also less intense than the parent support, accordingly the a0 is of 9.3 ±0.1 nm for all the bi-metallic materials (Table S1) [15]. Wide angle XRD registered for calcined samples are presented in Figure 2. Reflections corresponding to CuCo5:0 material (figure 2-e It is interesting to note that a slight decrease in crystallite size of the spinel phase is observed when decreasing the amount of copper (from 9.0 nm (CuCo4:1) to 8.1 nm (CuCo1:1)). For CuCo4:1 small crystallites of 7 nm for Co3O4 are obtained but very large CuO particles are formed with a size of 29 nm, as evidenced by the presence of very sharp reflection lines. Such results confirm that CuO particles are located outside of the SBA-15 mesoporosity. Representative TEM and HAADF analyses together with the EDX analysis were performed to provide deeper insight about the morphology, the dispersion and the chemical composition of the Cu and Co elements for freshly calcined materials (Figure 3, 4, S2, S3 and S4). First, monometallic supported oxides (CuCo5:0 and CuCo 0:5) exhibit different features with respect to their mixed oxide counterparts. As seen in figure S2, a distribution of polycrystalline cobalt oxide particles is clearly noticed as nanorods inside the SBA-15 mesoporosity for the CuCo 0:5 sample, while the single copper sample is showing a majority of large aggregates located outside the support porosity. Local EDX analyses were realized on selected regions from TEM images and are illustrated in Figure S3. The large particles found outside the SBA-15 are assigned to copper oxide particles (S3 1,3). Now, when analysing the distribution of copper oxide nanoparticles for CuCo4:1, CuCo1:1 and CuCo1:4 samples, we noticed that higher the amount of copper is, the larger particles are found (> 50 nm) outside the support porosity of silica (Figure 3). The main reason of such observations can be assigned to the high mobility of copper particles rendering its stabilization difficult on a silica type support because of the weak metal-support interactions [16]. On the contrary, CuCo1:4 exhibits no visible CuO NPs outside the mesoporosity (see figure 3-e), suggesting the successful insertion of copper species within the spinel structure. It is possible to find cobalt either in pure Co3O4 or in partially substituted CuxCo(1-x)Co2O3 spinel phase (see figure S3 andS4). Local analysis by EDX is confirming XRD patterns with the presence of the spinel oxides containing copper which are visible as confined NPs within the mesoporosity of the silica (figure 3e,f andS4).

Hence, it is interesting to note that the presence of cobalt is improving the insertion of copper within the spinel phase. On the other side, the dispersion of cobalt oxide NPs throughout the porous matrix is visibly improved by decreasing progressively the copper ratio. (Fig. 3 (e,f)). As observed, the NPs size is slightly decreased from 10 nm for pure Co3O4 to 6.8 nm for CuCo1:4 (Table S1). Complementary results were collected on three unique particles with different Cu-Co ratios (figure S4) from (HR)TEM, Fast Fourier transformation (FFT), EDX analysis and ICP-EOS. As observed by FFT/EDX, the plans associated to the spinel oxide are visible for CuCo 4:1, CuCo 1:1 and CuCo 1:4. Such plans allowed us to depict the 3D illustration of the planar disposition of the atoms (figure S4, FFT). Hence, the FFT data is associated to the typical cubic geometry of the Fd3m space group for all the three particles. However, the planar distances corresponding to the CuCo or Co3O4 spinel give similar values which avoid us to clearly identify the localisation or the distortions eventually generated by the enclosure of copper species in the tetrahedral sites of Co 2+ . Finally, ICP analysis of the materials (presented in the Table S1 and figure S4) is showing a good match between experimental values and the theoretical ones corresponding to the quantity of transition metal used during the preparation. Nevertheless, the EDX analysis on the single particles (figure S4) is indicating atomic percentages of cobalt from 57 at.% to 88 at.% and for copper from 43 at.% to 12 at.%. Knowing that the percentage of copper for the full substitution within CuxCo(1-x)Co2O3 is of 12.5 atoms, we can presume that all Co 2+ sites are substituted by the copper species in CuCo1:4 sample (figure S4) [25][26][27][28]. In the case of CuCo1:1 and CuCo4:1, knowing that Cu 2+ species are exceeding 12 at.%, the copper is favoured to be found as copper oxide decorating the surface of the spinel. To have more precise and local information about the copper and cobalt dispersion within the spinel and over the SBA-15 support, elemental mapping-EDX was realized on the calcined solids and the results are showed in figure 4. For all ratios, the confined NPs are composed of both Cu and Co, confirming again the formation of the CuxCo(1x)Co2O3 spinel phase. Besides, it is interesting to observe for all samples a population of small copper NPs (figure 4b,e,h) well distributed throughout the surface of the silica.

Reducibility properties

H2-TPR profiles were monitored in order to probe the temperature reduction domains of the different supported mono-and bi-metallic materials (figure 5). For the Co3O4 sample (CuCo 0:5), a classical two-step reduction process takes place with the occurrence of two distinct consumption ranges between 200 and 350 °C (Co(III) to Co(II) reduction) and between 350°C and 650°C (Co(II) to Co(0) reduction). The second reduction domain can be assigned to large particles situated at the external surface of the SBA-15 support, poorly dispersed, with a maximum reduction peak at 540 °C attributed to the presence of confined Co3O4 NPs [24,37]. On the other side, the profile reduction of monometallic copper oxide presents two peaks at 284 °C and 380 °C respectively, both confirming the presence of large aggregates of CuO clusters. [15] As observed in the case of our previous study based on CuNi systems [15], an intermediate situation is found for the bimetallic CuCo materials. Albeit a main reduction peak, located at ~ 307 °C, can be attributed to the simultaneous reduction of cobalt cations from (CuxCo(1-x)Co2O4) spinel phase, a clear shoulder is visible at lower temperature domains between 200 °C and 230 °C and is related to the reduction of well dispersed copper inserted into the spinel phases as previously observed by XRD and TEM (Figures 2 and4). We can also observe a small reduction domain at higher temperature (between 500°C and 600°C), suggesting the final decomposition step of spinel phases via the reduction of Co(II) in Co(0) species.

The evolution of crystalline phases and their thermal stability during the temperature reduction was monitored by in-situ DRX (figure 6 andS5). It can be seen that, in agreement with the chemical composition of the samples, the X-Ray diagrams display the peaks assigned to various crystalline phases of Cu-Co solids. For example, for the mono-metallic cobalt (figure S5), the diffraction pattern is showing the presence of Co3O4 exposed up to 350 °C, then the formation of CoO phase could be observed even after reduction at 600 °C. Small diffraction peaks associated to the metallic phase of Co, located at 51.5 ° and 75 °, can be observed upon reduction above 550°C (figure S5). Its intensity increases slightly accordingly to the increase of the reduction temperature, however the intensity signal is too weak to apply the Scherrer equation, suggesting the formation of small NPs. for Co 2p 1/2 on the pre-treated samples. The presence of two broad shake-up satellites is also evidenced and confirms the signature of Co 2+ / Co 3+ species from spinel Co3O4 phase. Upon reduction from 200 °C to 500 °C, the main photopeaks are shifted to 776.6 eV and 777.0 eV respectively, for both bimetallic samples, suggesting that the reduction is inducing the formation of the metallic cobalt species [39]. After reduction, the spectral information corresponding to Si 2p is showing the binding energy of 102.3 eV to 103.0 eV, corresponding to the presence of SiO2, notably from SBA-15 materials, consequently the binding energies of O 1s (532.05 eV and 533.4 eV) is associated to SiO2 support [40].

Catalytic reactivity toward cinnamaldehyde hydrogenation reaction.

The catalytic materials, obtained after appropriate reduction under hydrogen, have been tested in the liquid phase hydrogenation reaction of cinnamaldehyde using two sets of reaction conditions: (i) atmospheric pressure with a temperature of 150 °C and (ii) 10 bars with a temperature of 130 °C. Thus, the results illustrate the influence of reaction conditions and the influence of the mass ratio between the two metals on the catalytic performances.

•

Results obtained for catalytic runs performed at atmospheric pressure and with a temperature of 150 o C Conversion and selectivities obtained for all catalysts are given in figure 11-A and table 1. It can be seen that the catalytic activity of mono-metallic copper-based materials is close to zero, which is in agreement with the different structure and morphology observed for CuO NPs in CuCo 5:0 sample with respect to the bi-metallic materials.

[37] For this reason, the results about selectivity obtained on this catalyst are not relevant and are not presented herein. The introduction of copper within the spinel phase clearly led to a dramatic improvement of the CNA conversion, with almost 100 mol % CNA conversion and a high selectivity toward CNOL (72 mol%) obtained for the catalyst Cu:Co1:4. These results are in accordance with the data collected from TPR, in-situ XRD and TEM on the reduced bimetallic materials, as it is considered that the decrease in catalytic activity is caused by the increase of copper particle size, especially for sample CuCo4:1. [37] In addition, since the activity of this catalyst is similar to that of mono-metallic cobalt catalyst, it is likely that the catalytic activity is mainly attributed to cobalt. Regarding now the selectivity, the hydrogenation of the C=O double bond of cinnamaldehyde seem to be favoured albeit there are no major influences of the composition and particle size on the catalytic selectivity. Thus, for sample CuCo0:5, the selectivity to CNOL at small conversions is around 80 % and it decreases gradually up to ~ 50 % with increase in conversion. Instead, for bimetallic catalysts CuCo1:4 and CuCo1:1, it can be observed a selectivity stabilized around 70 %, up to a conversion of ~ 40 %, after that it suddenly decreases. This behaviour can be explained by the synergistic action of copper and cobalt, the latter being responsible for the increased selectivity to CNOL. [4] As reported in literature [3,4], cobalt is known to be selective for CNOL due to the width of the band d, while it is maintained to a constant value for copper. Indeed, the increasing of copper amount in the sample (CuCo4:1) decreases the selectivity towards 60 %, but its evolution with conversion follows the trend observed for the two bimetallic samples.

Table 1. Results of catalytic activities as well as the selectivities measured for the materials presented in this work.

Figure 1 .

 1 Figure 1. a) Adsorption/Desorption isotherms; b) Pore size diameter corresponding to support and Cu/Co containing materials: a) SBA-15; b) CuCo0:5, c) CuCo1:4; d) CuCo1:1; e) CuCo4:1, f) CuCo5:0.

  ) point out the presence of large copper oxide NPs (ICDD 48-1548), with a mean crystallite size around 32 nm evidencing a strong NPs enrichment on the external surface of SBA-15 granules. When switching from CuCo 5:0 to CuCo 0:5 solid, a significant mitigation of the reflexion lines occurs, suggesting a better distribution of cobalt oxide NPs within the mesoporosity of the silica. Co3O4 phase was further identified by comparison with the reference (ICDD 42-1467) and NPs size was estimated around 9 nm. For mixed CuCo ratios, the observed diffraction patterns did not allow us to discriminate Co3O4 phase from CuxCo(1-x)Co2O3 phase, [36] which are both exhibiting very close 2θ values. Besides, the CuCo bimetallic solids present diffraction patterns which can be assigned to CuO and spinel phases, all displaying different relative intensities in agreement with the ratio of Cu and Co requested for the preparation (figure 2).

Figure 2 .

 2 Figure 2. Wide angle XRD patterns for: (a) CuCo0:5; (b) CuCo1:4; (c) CuCo1:1; (d) CuCo4:1; (e) CuCo5:0; bottom: in blue CuO (pdf No: 48-1548) and in red Co3O4 (pdf No: 42-1467).

Figure 3 .

 3 Figure 3. TEM images corresponding to: (a,b) CuCo4:1; (c,d) CuCo1:1; (e,f) CuCo1:4.

Figure 4 .

 4 Figure 4. (a,d,g) Representative HAADF images obtained for calcined CuCo based materials; (b,e,h) copper localization as determined by STEM-EDX mapping analysis; (c,f,i) cobalt localization as determined by STEM-EDX mapping analysis.

Figure 5 .

 5 Figure 5. TPR profiles for: (a) CuCo0:5; (b) CuCo1:4; (c) CuCo1:1; (d) CuCo4:1; (e) CuCo5:0.

Figure 9 .Figure 10 .

 910 Figure 9. Core level spectra of: Co 2p (a), Cu 2p (b) Cu LMM (c) obtained for CuCo1:1 ( in situ treatment under air/500°C, H2/200 °C, H2/300 °C, H2/500 °C

Figure 11 .

 11 Figure 11. Catalytic activity for: (a) CuCo0:5, b) CuCo 1: 4; (b) CuCo 1:1; (d) CuCo4:1, e) CuCo5:0 ;performed at: A) Reaction conditions:Tred = 500 °C for sample CuCo0:5 and 350 °C for the others; Treaction = 150 °C, 0.265 g catalyst; 1 mL CNA, 25 mL propylene carbonate as solvent, flow rate of H2 = 1 L h -1 , speed of agitation = 900 rpm). B): Tred = 500 °C for sample CuCo0:5 and 350 °C for other samples; Treaction = 130 °C, 0.133 g catalyst; 1 mL CNA, 40 mL isopropanol as solvent, agitation rate = 750 rpm.

For CuCo1:4 and CuCo1:1 samples, it can be seen that the peak corresponding to the Cu 0 phase (2θ ~ 50.2 °) is already identified at 350 °C and is relatively well-defined. However, its intensity is rather constant throughout the temperature reduction range from 350 °C up to 750 °C, indicating the successful stabilization of copper onto the SBA-15 support when cobalt is added. It is interesting to note that the XRD patterns of CuCo1:4 recorded with a reduction temperature of 350 °C show weak diffraction peaks corresponding to Cu 0 and Co 0 [24]. Taking into account the TPR profile for this sample, a significant shoulder is observed at ~200 °C and is likely due to the reduction of smaller copper species from nanoparticles of (CuxCo(1-x)Co2O4) spinel, which are not detected by insitu XRD at 250°C. For CuCo4:1, the diffraction pattern is indicating very intense peaks for CuO crystalline phases, in well agreement with the TEM images. The cobalt oxide peaks are having very low intensity, being visible only for temperatures lower than 250 °C. The presence of H2 is revealing the metallic copper at temperatures higher than 350 °C by the appearance of characteristic peaks. Moreover, any peaks corresponding to metallic cobalt cannot be identified in XRD patterns, indicating their high dispersion. A ratio of 1:1 between the two elements (CuCo1:1) leads to a slight increase in the amount of copper that is not involved in spinel phase. For the pure cobalt-based sample, similar diffraction peaks were found with respect to CuCo1:4.

Effective localization and elemental composition of the NPs, after their progressive reduction in H2 up to 750 °C, were registered by TEM and HAADF analyses for all bimetallic materials. A complementary EDX analysis was used in order to have local information about the analytical composition. Selected images corresponding to the Cu-Co materials are illustrated in Figure 7 and 8 and the corresponding EDX analysis are found in Figure S6. For all bimetallic materials, the presence of confined NPs can be observed after reduction under hydrogen (Figure 7).

The size of the metallic NPs is close to its oxidic counterparts, indicating a high thermal stabilisation even after reduction at 750°C. On the rich cobalt sample (CuCo 1:4), EDX analysis indicate the presence of the two metals in the same nanoparticle, with a mass ratio close to the theoretical one (figure S6, EDX 3). This result supports the hypothesis of heterostructures formation for CuCo 1:4 ratio. Albeit NPs contain as well both metals, different features were found in CuCo1:1 sample since different mass ratios were obtained, with experimental values being close to the theoretical one or close to copper-rich compositions, in particular in the case of nanoparticles located on the external surface. However, cumulating the TPR and TEM results, these extra-porous nanoparticles constitute a minor population. In Figure 8 are presented HAADF images as well as STEM-EDX mapping on the selected images. The results are showing that the confined NPs are maintaining a similar ratio with respect to the calcined materials. It is interesting to note that a high dispersion of copper species is present onto the silica, especially for the ratio of CuCo1:1 where the appearance of small copper NPs is revealed after the reduction. In the case of CuCo1:4, small copper based NPs are homogeneously distributed throughout the support porosity, suggesting a high dispersion which arises from a high thermal stability even after treatment in H2 up to 750 °C. Such results confirm that the presence of cobalt is critical toward the stability of finely dispersed Cu species initially located inside the spinel structure. Moreover, the presence of cobalt is suggested to favour the formation of highly dispersed CuO NPs. Complementary surface information is given in Figure 9 and 10 and the spectral information from XPS are listed in Table S2 for CuCo1:4 and CuCo1:1. The materials were pre-treated under air at 500 °C prior to the temperature programmed reduction under hydrogen. The XPS spectra for Cu 2p (figure 9 and 10) registered before the reduction treatment shows the typical shake-up satellite corresponding to CuO species for the pre-treated solid, consequently the presence of the Cu 2+ state in octahedral position is confirmed by the peak at 934.7 eV and 933.2 eV for both CuCo1:4 and CuCo 1:1. While reducing the samples at 200°C, the shake up satellite is disappearing and the main peak is becoming sharper. Hence, Cu 2p photopic is progressively shifted to 931.9 eV, which is typicaly associated with the presence of metallic copper species [38]. Interestingly, similar catalytic performances were obtained for the same ratio in the case of catalysts based on CuNi/SBA-15 as reported in a previous work [15]. These results suggest that (i) a ratio M1: M2 = 1:4, where M1 = Cu and M2 = Ni(Co) represents the optimum ratio at which the synergistic interactions between the two metals have positive effects on the physico-chemical and catalytic properties of materials and (ii) as also shown for the CuNi system, the bimetallic particles obtained by our approach are not classical alloys, but rather hetero-structures which share a mixed interface, with no segregation of copper on the surface of cobalt or nickel. For CuCo system, the results indicate similar catalytic behaviours as CuNi system. Additional increase of the copper amount

Sample

Acknowledgments

The authors would like to thank Pardis Simon (UCCS) for her help in carrying out the work on the Surface Analysis facilities. The Chevreul Institute is thanked for its help in the development of this work through the ARCHI-CM project supported by the "Ministère de l'Enseignement Supérieur de la Recherche et de l'Innovation", the region "Hauts-de-France", the ERDF program of the European Union and the "Métropole Européenne de Lille.


Melodie Galerne

Benoît Heinrich

Arnaud Hemmerle

Philippe Fontaine

Nicolas Giuseppone
email: giuseppone@unistra.fr

Patrick Kékicheff
email: patrick.kekicheff@ics-cnrs.unistra.fr

Spatially Addressed Supramolecular Nanowires: A Full Structural Characterization by GIWAXS

Keywords: supramolecular polymers, nanowires, supramolecular electronics, orientational alignment, microelectrodes, grazing incidence wide-angle X-ray scattering

Supramolecular nanowiring is based on the precise interconnection of nano-or microelectronic circuit elements with electroactive supramolecular polymers displaying directional charge transport properties. It is a particularly attractive topic of current research in order to access miniaturized organic electronic devices. The bottom-up construction of such integrated nanowires can be generated by specific supramolecular polymerization processes -based on addressed nucleation growth mechanisms -which provide a spatial precision in their placement that cannot be accessed by top-down approaches. However, to implement, quantify, and validate such axis. From a broader perspective, this work highlights the high potential of GIWAXS to become a reference method for achieving the full characterization of nanowiring processes with various types of supramolecular polymers and polymer self-assemblies integrated in complex device geometries.

supramolecular interconnects in nanotechnologies, powerful characterization techniques, which probe both the fine structural and orientational features of the nanowires, are still required. Most techniques usually lack the ability to access the morphological parameters of the targeted nanostructures lodged within the electronic circuits at that scale (for instance between micro/nanoelectrodes). Even grazing-incidence wide-angle X-ray scattering (GIWAXS) poses major hurdles for the practical investigation of such miniaturized electronic devices due to beam path obstruction for most in-plane orientation of the samples. To overcome these limitations, we have implemented a sensing modality based on GIWAXS collection for unique sample orientation and spot shape analysis in reciprocal space allowing the entire film morphology to be accessed. In this case study involving triarylamine-based nanowires, we fully determine in a single experiment the packing structure of the supramolecular polymers interconnecting microelectrodes, together with their 3D and 2D orientational order parameters (S3D ≈ 0.986 and S2D ≈ 0.91). Our investigation demonstrates that the nanowires lie not only flat along the substrate but are also aligned along the normal to the electrodes. Furthermore, this alignment occurs with a coherent length of more than 100 nm, representing more than 200 molecular lengths along the nanowire

INTRODUCTION

One of the most challenging objectives in supramolecular electronics [START_REF] Schenning | Supramolecular electronics; nanowires from selfassembled p-conjugated systems[END_REF][START_REF] Su | The Art of Aligning One-Dimensional (1D) Nanostructures[END_REF][START_REF] Jain | New directions in supramolecular electronics[END_REF] is to precisely interconnect circuit elements, such as nano-and micro-electrodes, with supramolecular organic nanowires. Because charge transport is highly directional, the entire device efficiency depends on both the (local) structural and (long range) orientational alignments of the organic (semi)conducting material. Hence, monodimensional supramolecular polymers with a high degree of internal organization (e.g. large number of aligned crystalline domains), and an optimal orientation between electrodes, are expected to considerably increase the charge transport efficiency and miniaturize organic electronic devices. [START_REF] Jain | New directions in supramolecular electronics[END_REF][START_REF] Moulin | Advances in Supramolecular Electronics -From Randomly Self-Assembled Nanostructures to Addressable Self-Organized Interconnects[END_REF] This has been established for e.g. OLEDs, OFETS, [START_REF] Jo | Large-Scale Alignment of Polymer Semiconductor Nanowires for Efficient Charge Transport via Controlled Evaporation of Confined Fluids[END_REF][START_REF] Chaudhary | Highly Aligned and Crystalline Poly(3-Hexylthiophene) Thin Films by off-Center Spin Coating for High Performance Organic Field-Effect Transistors[END_REF][START_REF] Chang | Macroscopic Alignment of One-Dimensional Conjugated Polymer Nanocrystallites for High-Mobility Organic Field-Effect Transistors[END_REF][START_REF] Xiang | Ge/Si Nanowire Heterostructures as High-Performance Field-Effect Transistors[END_REF] solar cells, [START_REF] Awartani | In-Plane Alignment in Organic Solar Cells to Probe the Morphological Dependence of Charge Recombination[END_REF][START_REF] Yang | Effects of Nanostructure Geometry on Polymer Chain Alignment and Device Performance in Nanoimprinted Polymer Solar Cell[END_REF][START_REF] Yang | Aligned Carbon Nanotube Sheets for the Electrodes of Organic Solar Cells[END_REF] and chemical sensors. [START_REF] Peng | Aligned Coaxial Nanowires of Carbon Nanotubes Partially Sheathed with Polyaniline for Chemical Sensors[END_REF][START_REF] Mcalpine | Highly Ordered Nanowire Arrays on Plastic Substrates for Ultrasensitive Flexible Chemical Sensors[END_REF] The long range anisotropy and orientational alignment of various kinds of supramolecular nanoobjects, such as fibers, rods, chains, and belts are most often studied by microscopy techniques, such as polarized optical microscopy, AFM, SEM, and TEM. [START_REF] Persson | In Situ Observation of Alignment Templating by Seed Crystals in Highly Anisotropic Polymer Transistors[END_REF] In addition, spectroscopy techniques such as UV-vis absorption, spectroscopic ellipsometry (SE), and near-edge X-ray absorption fine structure (NEXAFS), are commonly used to probe the local organization of fibrillar structures. [START_REF] Persson | In Situ Observation of Alignment Templating by Seed Crystals in Highly Anisotropic Polymer Transistors[END_REF][START_REF] Nahid | NEXAFS Spectroscopy of Conjugated Polymers[END_REF] However, the precise molecular order of the supramolecular polymer backbone, from which the (nano)object is built-up, can only be probed and quantified by more local characterization techniques such as x-ray scattering. Though, when only a few supramolecular nanowires are arranged in an electronic circuit, the scattering experiment becomes more complex because of the low signal intensity and device geometry. We reasoned that, in principle, these limitations can be overcome by using grazing incidence wide-angle X-ray scattering (GIWAXS), as large areas within thin surficial layers are probed, with the advantage to elucidate the local molecular packing structure of the nanowires. [START_REF] Zarubin | Improved Anisotropic Thermoelectric Behavior of Poly(3,4-Ethylenedioxythiophene):Poly(Styrenesulfonate) via Magnetophoresis[END_REF] Moreover, by measuring the sample in several directions, one can also potentially clarify the alignment and orientation of the polymer in their effective wiring configuration. [START_REF] Lin | Nonvolatile Organic Field Effect Transistor Memory Devices Using One-Dimensional Aligned Electrospun Nanofiber Channels of Semiconducting Polymers[END_REF][START_REF] Bastianini | In-situ Monitoring Poly(3-Hexylthiophene) Nanowire Formation and Shape Evolution in Solution via Small Angle Neutron Scattering[END_REF] Finally, the coherence length of the supramolecular polymers may be obtained, [START_REF] Weller | Highly Efficient and Balanced Charge Transport in Thieno[3,4-c] Pyrrole-4,6-Dione Copolymers: Dramatic Influence of Thieno[3,2-b] Thiophene Comonomer on Alignment and Charge Transport[END_REF] and a GIWAXS derived orientation parameter calculated. [START_REF] Gujral | Structural Characterization of Vapor-Deposited Glasses of an Organic Hole Transport Material with X-Ray Scattering[END_REF] Therefore, one can expect that a single GIWAXS experiment could ideally provide a fully quantitative multiscale structural and orientational characterization of the supramolecular wires directly measured in the organic electronic device. Nevertheless, to the best of our knowledge, the necessary developments to implement such a powerful characterization method in the context of nanowired devices have not been investigated in the literature so far.

As a proof of principle, we decided to study the local structure and long-range anisotropy of self-assembled triarylamine-based supramolecular polymers [START_REF] Moulin | Triarylamyne-Based Supramolecular Polymers: Structure, Dynamics and Functions[END_REF] placed between microelectrodes (Fig. 1). This new family of supramolecular polymers has been previously shown to display excellent electronic and photonic conducting properties. [START_REF] Faramarzi | Light-Triggered Self-Construction of Supramolecular Organic Nanowires as Metallic Interconnects[END_REF][START_REF] Armao | Long-Range Energy Transport via Plasmonic Propagation in a Supramolecular Organic Waveguide[END_REF][START_REF] Armao | Supramolecular Organic Nanowires as Plasmonic Interconnects[END_REF][START_REF] Busseron | Light-Controlled Morphologies of Self-Assembled Triarylamine-Fullerene Conjugates[END_REF] Their local supramolecular structures were characterized by various techniques, including powder X-ray diffraction (XRD). [START_REF] Moulin | Triarylamyne-Based Supramolecular Polymers: Structure, Dynamics and Functions[END_REF] Moreover, we recently demonstrated the possibility to grow triarylamine (TAA) nanowires directly in between electrodes by introducing a so-called supramolecular electropolymerization process starting from TAA 1 (Fig. 1). [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] The spatially addressed and oriented formation of these nanowires is accessible by a multistep oxidation / nucleation / growth mechanism (Fig. 1b-d). First, a solution of TAA 1 in heptane is drop casted onto interdigitated electrodes, while applying a direct current electric field of 1.25 × 10 7 V•m -1 . Neutral TAA 1 molecules, that come in direct contact with the anode, are electro-oxidized into their radical cation form (TAA 1 •+ ), subsequently triggering their nucleation by charge transfer complexation and hydrogen bonding. The anode thus becomes a seeding surface from which the supramolecular polymers can grow: neutral TAA 1 stack onto the nuclei and the nanowires polymerize from the anode towards the cathode, following the lines of the applied electric field. Finally, the supramolecular polymers laterally aggregate with one another by van der Waals forces to produce larger bundles of supramolecular nanowires. This novel supramolecular electropolymerization process warrants unprecedented control over the localization of the interconnects, by (i) addressing electro-oxidation of the TAA 1 monomers at the anode, and by (ii) inducing directional growth of the supramolecular polymers in the applied electric field. The 2D orientational alignment of the supramolecular polymers was quantified for the direction normal to the electrodes, using SEM and AFM images, with order parameters of 0.87±0.09 (at a concentration of 2 mg/mL) and 0.94±0.08 (at 5 mg/mL). [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] However, the 3D orientational parameter, as well as the local supramolecular structure and its coherence length are not accessible by these techniques but of first importance in the context of a nanowiring process.

We now present a much more powerful analysis by making use of a single GIWAXS experiment to fully characterize this system (Fig. 1e). The main challenge arises from the fact that a method, capable of delivering the local information within the electronic circuit, by directly characterizing the nanowires in between microelectrodes, is required in order to take optimal advantage of our systems configuration. Classical methods for deriving orientational order parameters, by means of rocking curves, cannot be used as the thin film has the same height (~100 nm) as the microelectrodes, that consequently obstruct the path of the grazing incident beam. Furthermore, inplane rotation of the microelectronic device would inevitably lead to inherent variations, both in the effective scanned volumes and in the background scattering contributions made by the glass substrate. To overcome these experimental limitations, we resorted to a direct analysis of the intensities of Bragg reflections in the GIWAXS pattern. Here, we combine the radial and angular broadening observed for each spot in reciprocal space collected for the same configuration (incident beam parallel to the microelectrodes, w = 0, Fig. 1e) to infer the morphological parameters of the film. With this technique, not only the (supra)molecular organization of the nanowires can be determined, but, through the acquisition of their coherence length and both 2D and 3D orientational parameters, also their long-range orientation. 26 with permission] formed by supramolecular electropolymerization of TAA 1 solutions in heptane at 2 mg/mL (c) and 5 mg/mL (d); the solutions were drop casted onto 2.5 mm long interdigitated gold micro-electrodes (width 10 µm; height 100 nm; spaced away by 10 µm) at a 1.25×10 7 V/m applied DC electric field [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] ; (e) Schematic of the GIWAXS setup on the SIRIUS beamline at the French National Synchrotron Radiation facility SOLEIL (Saint Aubin, France) [START_REF] Fontaine | Soft interfaces and resonant investigation on undulator source: A surface X-ray scattering beamline to study organic molecular films at the SOLEIL synchrotron[END_REF] including relevant geometry definitions. The sample is tilted by 0.12° with respect to the incident beam for grazing incidence wide-angle measurements. The sample can be rotated about its vertical axis: its position is defined by the angle w made by the normal to the microelectrodes with respect to the normal to the incident beam. In configuration where the incident beam is parallel to the microelectrodes (w = 0), all the electrode gaps, which are filled with TAA 1 nanowires, are irradiated simultaneously, as the x-ray footprint is 500 µm wide. The angles f and b define the disorientation of the long axis of a nanowire with respect to the substrate plane and to the normal to the electrodes respectively. The horizontal ("longitude") angle, µ, and the vertical ("latitude") angle, d, define the direction of the diffracted beam with respect to the incident beam. The diffracted beam hits the 2D detector (detector rotation not shown) with a total scattering angle 2q with q the Bragg angle.

RESULTS AND DISCUSSION

Structure of the supramolecular nanowires films. The GIWAXS patterns, collected at low (2.5 mg/mL) and high concentration (5 mg/mL) of TAA 1, once the electropolymerization was processed, are each composed of up to 40 reflections, going from 17 to 3 Å in real space (see Fig. 2 and Fig. S1 in ESI †). Through the examination of the positions of the diffraction spots and the averaging of the radial scattering, the structure of the supramolecular polymer film is resolved.

Both concentrations lead to the same structural determination. The low-angle region is composed of (i) two intense first order reflections (0.3525 Å -1 and 0.3726 Å -1 ) and (ii) a series of weaker higher order reflections of (hk0) type, occurring from a bidimensional sub-lattice of rectangular symmetry. The long and short edges of this sub-lattice are assigned to the a and b-axes, with a = 33.7(3) Å and b = 21.0(6) Å. The systematic missing (hk0) peaks with odd h + k sum suggests a centered symmetry for the sub-lattice. [START_REF]International tables for Crystallography[END_REF] The wide-angle region includes crossed reflections involving the periodicity along the c-axis, from which the three-dimensional structure is determined. The indexation converges on a crystal lattice of monoclinic symmetry. A simple centered 3D lattice cannot be brought into agreement with the measured sets of q values, since some (hkl) peaks with odd h + k + l sum can be detected. Thus, the peak positions of all samples can be indexed to a monoclinic lattice according to:

𝑞 hkl = 2𝜋 %& ℎ 𝑎 ) * + & 𝑘 𝑏 sin𝛼 ) * + & 𝑙 𝑐 sin𝛼 ) * -2 𝑘𝑙 𝑏𝑐 cos𝛼 sin * 𝛼 3 4/* (1)
with 3D lattice constants, a = 33. The volume of the unit cell is 6870 Å 3 and contains Z=8 molecules. Upon supramolecular electropolymerization, TAA 1 forms a crystalline structure of equivalent symmetry and a geometry that is close to the geometry of previously reported triarylamine monoamide molecules that selfassemble in solution after photooxidation. [START_REF] Nyrkova | Supramolecular Self-Assembly and Radical Kinetics in Conducting Self-Replicating Nanowires[END_REF] It is clearly evidenced here that supramolecular polymerization preserves the characteristic molecular shape and self-assembly by twinning of aryl segments through H-bonds into double-strands surrounded by a shell of aliphatic chains. Doublestrands then arrange at the nodes of rectangular columnar a × b sublattice with somewhat shrunken a/b ratios, as compared to hexagonal geometry (a/b = √3), and with a periodicity of two molecular thicknesses along the c-axis. Non-orthogonal angle a could here be determined, implying that the double-strands lie along the b-axis with a longitudinal shift of 1.6-1.7 Å.

A schematic representation of the molecular organization can be proposed from the GIWAXS analysis (Fig. 3). When piling, TAA 1 molecules present a colinear arrangement of their central nitrogen atoms separated by the distance hmol of 0.5 nm. Because of the propeller shape of the TAA units, their aromatic rings come at closer stacking contact which also enforces them to alternate their helical chirality between D and L enantiomers. Because of the 60° tilt of TAA along their elongation axis, the characteristic repeating unit h is equal to 1 nm (2 x hmol). Each monocolumn of TAA is laterally associated to another one by an infinite string of hydrogen bonds provided by the amide groups. contributes to the scattering. However, this information is difficult to extract because it is hindered by a strong scattered background signal, due to the glass substrate in the range 1.2 to 1.6 Å -1 (see ESI †). Indeed, at such incidence angle, the beam foot-print is about 50 mm long, whereas the effective supramolecular polymer film length is 450 µm in this geometry (sample composed of 15 electrode pairs of 2.5 mm long, having each a width of 10 µm and spaced away by 10 µm). Thereby, in the following, our focus is mainly on the configuration where the incident x-ray beam is parallel to the microelectrodes (w=0). In this geometry, the scattering patterns arise from the bulk of the supramolecular polymers grown in the 30 gaps between the 15 parallel interdigitated gold microelectrodes. With the incident beam parallel to the electrodes, the scattering vector q probes diffraction perpendicular to the nanowires, both in the in-plane and out-of-plane directions. In the presented data, a coordinate system is defined for the scattering vector q associated with the GIWAXS patterns, with qz defined as the component of q in the specular direction (perpendicular to the silica substrate) and qy defined as the q component parallel to the substrate.

Whatever the concentration, the (002) diffraction peak appears as a spot located on the qy axis in the configuration where the incident x-ray beam is parallel to the microelectrodes (Fig. 2a and Fig. S1). This indicates that the nanowires are strongly confined to an in-plane configuration. As it will be shown below, the sharpness of this peak suggests that the nanowires lie not only flat along this axis, but are also well aligned along the normal to the electrodes. These conclusions are further advocated by the absence of (002) spot when the samples are rotated by 90° (Fig. S1), beyond the experimental limitations already mentioned. The c-axis is thus chosen to be oriented along the long-axis of the nanowires and the rectangular sublattice a x b is therefore in a plane almost orthogonal to the substrate (the angle a between the b-axis and the c-axis being ~94.(5)°).

At low concentration, the arc-shaped reflections, observed in the GIWAXS pattern for all the out-of-plane directions (Fig. 2a), indicate a rotational freedom of the crystal along its long c-axis in comparison with the lamellar stacking vector. However, at this concentration, the rings are not fully isotropic and the signal concentrates around the peak positions with limited angular distributions of the arcs. This suggests the domination of preponderant structural motifs with a "roll" of the crystals remaining moderate. Remarkably, at this concentration of nanowires, the GIWAXS pattern arises mainly from two preferential orientations of the crystals (Fig. The planes parallel to the columns faces are less dense because they are packed in a staggered configuration. The fact that anisotropic objects tend to align their leading crystallographic planes along substrates, has been observed in many systems, including lyotropic liquid crystals which undergo topological transformations from cylinders to lamellae encompassing different structures, such as anisotropic ribbons and connected rods. [START_REF] Kékicheff | Crystallography of Systems with Long Periods: a Neutron-Scattering Study of Sodium Dodecyl Sulfate/Water Mesophases[END_REF][START_REF] Kékicheff | From cylinders to bilayers: A structural study of phase transformations in a lyotropic liquid crystal[END_REF][START_REF] Mcgrath | Spiral textures in lyotropic liquid crystals: first order transition between normal hexagonal and lamellar gel phases[END_REF] Note that these preferential orientations are lost when the nanowires are grown in a more concentrated solution. Indeed, at higher concentrations almost uniquely isotropic rings are observed in the out-of-plane directions. The a* and b*-axis now have random orientations, even though the c*-axis remains parallel to the substrate and perpendicular to the electrode direction, Fig. S1, ESI †). In contrast, at the lower concentration, arc-shaped reflections are discerned (Fig. 2a). This observation can be interpreted twofold. At high concentration, the tighter molecular packing reduces the flexibility gradient along the growth direction of the nanowires and preferential orientations of the rectangular sublattice a x b are not favored at the interface with the substrate. Indeed, steric restrictions may manifest themselves as freedom for the nuclei to place with no special orientation. Thereby, the crystals will subsequently grow along their long c-axis with no preferential orientations. Additional considerations can be suggested. In particular, the loss of out-of-plane spots, with local reinforced intensity in their limited angular arc-shaped reflections, may be apparent only. Therefore, it is possible that some preferential orientations are still present in the vicinity of the substrate interface at high concentration, but the signal coming from these orientations is dominated by the bulk of the more concentrated supramolecular film.

Parallelism of the nanowires to the substrate. In the configuration where the incident beam is parallel to the microelectrodes (w=0, Fig. 1e), the GIWAXS patterns share the same (002) diffraction peak, located along the in-plane qy-axis (Fig. 2a and Fig. S1, ESI †) for all concentrations.

This spot is similar in position to the one found in the (00l) progression with even Miller indices, characteristic of the stacking of the aromatic rings composing the triarylamine molecules in a centro-symmetric unit cell. [START_REF] Nyrkova | Supramolecular Self-Assembly and Radical Kinetics in Conducting Self-Replicating Nanowires[END_REF] Here, for the TAA 1 disk-like core motif, the d-spacing is 9.7 Å. The localization of the (002) spot on the qy-axis indicates that the nanowires are strongly confined to an in-plane configuration. Its sharpness suggests that the crystals lie flat along this axis with little out-of-plane tilt. The angular width of this strong equatorial reflection allows us to determine the degree of orientation of crystallites along the c-axis. In order to estimate the fraction of nanowires that is aligned and the degree of alignment, the GIWAXS intensity is integrated along the 2q axis and plotted as intensity I versus azimuth f.

Figure 4 shows the profile for the (002) spot measured at a concentration of 2.5 mg/mL (see Fig. 2). Gaussian and Lorentzian shape function can be differentiated by comparing the fits of the tail of the peak. A single Gaussian function does not reproduce the weak broad tail at the bottom of the Bragg singularity. Good fits of the profile are obtained using either the sum of two Gaussian functions or a Lorentzian function. As the value of the maximum intensity is uncertain at the precise equator location in Fig. 2, a small uncertainty arises in the width of the profile. Nevertheless, all the fitting functions lead to a similar full width at half maximum of the I002(F) profile: the FWHM is 3.3 ± 0.2 °, which means that the aligned crystals have a very small variation of their tilt angle with respect to the substrate.

Figure 4.

Azimuthal profile of the diffraction spot (002) after subtraction of the scattered background intensity, for the configuration in which the incident x-ray beam is oriented parallel to the micro-electrodes, and a concentration of 2.5 mg/mL is used (see Fig. 2). The profile (solid circles) can be fitted using either a Lorentzian or the sum of two Gaussian functions. They all lead to a similar full width at half maximum (FWHM) of 3.3 ± 0.2 °. The uncertainty is estimated from the unknown maximum value that the intensity would take at the equator (the angle f at exactly 90° is out-of-reach in the GIWAXS pattern).

The small spread of these lattice planes in the narrow out-of-plane distribution of the stacking peak about the surface normal (mosaic angle ~1.6°) can be further quantified by calculating an orientational order parameter. For orientation distributions that are centrosymmetric as usually encountered for liquid crystals and polymer fibrils, the first non-vanishing moment of the angle f between an element and the director is quadratic. The most common method to define the orientation of assemblies from x-ray scattering experiments is by calculating Herman's orientation parameter: 33

𝑆 3D = 1 2 (3〈cos * 𝜙〉 -1) (2)
where the mean-square cosine is calculated from the scattered intensity 𝐼(𝜑) by integrating over the azimuthal angle 𝜑 :

〈cos * 𝜙〉 = ∫ 𝐼(𝜑) cos * 𝜑 sin𝜑 𝑑𝜑 BC C ∫ 𝐼(𝜑) sin𝜑 𝑑𝜑 BC C (3)
For perfect horizontal orientation S3D = 1, while for isotropic orientation S3D = 0. Herman's orientation parameter refers to the second-order spherical harmonics of Legendre polynomial 𝑃 * (𝜙) of the cosine of that angle. Higher moments of the orientation distribution can also be obtained from the azimuthal profile of the reflection, and hence the full orientation distribution can be calculated, since the function can be expanded in series of even-order Legendre polynomials, [START_REF] Lowell | Molecular orientation distribution derived from an arbitrary reflection[END_REF] with

〈𝑃 *E (𝜙)〉 = ∫ 𝐼(𝜑) 𝑃 *E (𝜑) sin𝜑 𝑑𝜑 BC C ∫ 𝐼(𝜑) sin𝜑 𝑑𝜑 BC C (4)
By applying equations 2-4 onto the azimuthal profile of the (002) spot shown in Fig. 4, one obtains: which is the second lowest and most important order parameter. These values demonstrate that the dominant feature of the GIWAXS patterns is a strong confinement of the supramolecular nanowires to a 2D plane with their long axis (i.e. the crystal c-axis) being parallel to the substrate: the nanowires lie flat on the substrate with negligible out-of-plane tilt. Taking advantage of this remarkable result, we will consider hereafter that nanowires are positioned in the lateral plane of the substrate only.

𝑆 3D = 〈𝑃 * (𝜙)〉 =
Alignment of the nanowires along the normal to the microelectrodes. In this section the twodimensional disorientation of the nanowires is investigated. As already mentioned, the orientational distribution cannot be obtained by single rocking curves along a vertical axis, since accurate measurements of the diffracted signal is hampered by both variations in the footprint area and in the intensity of the grazing incident x-ray beam due to path obstruction by the microelectrodes that have the same height as the film (~100 nm height). To overcome these limitations, we resorted to the direct analysis of the intensities of Bragg reflections in the GIWAXS pattern. This approach makes use of the important constraint that the shapes of reflections at different positions on the detector are related for a given configuration (here, incident beam parallel to the microelectrodes). [START_REF] Millane | Relating reflection boundaries in X-ray fiber diffraction patterns to specimen morphology and their use for intensity measurement[END_REF] Indeed, the main sources of the finite size of the Bragg reflections are disorientation and finite crystallite sizes (in our synchrotron experiments the beam spread is negligible). The effects of finite crystallite size are directly related to the size of the reflection in reciprocal space while the effect of disorientation is to arc each reflection about the center of the diffraction pattern, forming part of a Debye-Scherrer ring.

As the growth of the nanowires follows the lines of the applied electric field, we assume that the disorientation of their long axis is axisymmetric. Each nanowire is made of molecular segments, packed together to form well-ordered crystallites within each object. the crystallites may have partial, weak, or no correlations in position and orientation, they nevertheless share the same common long axis of the nanowire in which they are embedded. This axis was defined conveniently as the z-axis in the real space. The c-lattice vector is the lattice vector parallel to that axis. In the following, the crystallites are considered to have a finite size and their average dimensions normal and parallel to the z-axis are noted 𝑙 lat (i.e. in the lateral direction) and 𝑙 axial (i.e. in the axial direction). The effects of finite crystallite size are directly related to the size of the reflection in reciprocal space (see also last section). The effect of disorientation is to rotate the reciprocal lattice point around the origin of reciprocal space, and the shape and width of the angular profile due to disorientation are approximately constant over most of reciprocal space.

If the substrate does not induce any preferential orientation, the crystallites are randomly

The preferential orientations of the sublattice a x b induced by the substrate makes the projection to be not cylindrically isotropic. Combining this with the small variation in planar inclinations of the nanowire long axis to the direction normal to the microelectrodes, the superimposition of these two effects has the effect of spreading out the diffracted intensity from reflection along Debye-Scherrer arcs. That is, the intensity at a given position of the detector is arced about the center of the pattern at constant diffraction angle. Following eqn.10 of the Methods section, and after corrections of the intensity and background subtraction, the profile of the i th reflection measured on the detector in polar coordinates (qi,Fi) can be written in the form

𝐺 S (𝑞, 𝜙) = 𝐼 S 𝐺 Z & 𝑞 -𝑞 S Δ𝑞 S ) 𝐺 \ & 𝜙 -𝜙 S Δ𝜙 S ) ( 6
)
where Ii is the peak intensity, Dqi and DFi are the radial and angular half-widths of the reflection of normalized profile shape Gq and GF. The variations of the half-widths, with position on the detector, have been calculated. [START_REF] Millane | Relating reflection boundaries in X-ray fiber diffraction patterns to specimen morphology and their use for intensity measurement[END_REF] On the detector, placed at a distance D from the sample rotated by an angle w to the normal of the incident beam, these half-widths read: where qi is the Bragg angle corresponding to the wavevector 𝑞 S = (4𝜋/𝜆) sin 𝜃 S in the reciprocal space at the wavelength l = 1.24 Å (energy 10 keV), and K a calibration factor transforming pixel sizes of the detector into Å -1 units. In these expressions, the constants u and v are related to the mean crystallite sizes normal and parallel to the nanowire axis (2u = 1/llat ; 2v = 1/laxial), and b to the mean disorientation.

Δ𝑞 S =
As noted by Millane, [START_REF] Millane | Relating reflection boundaries in X-ray fiber diffraction patterns to specimen morphology and their use for intensity measurement[END_REF] the precise meaning of "mean" depends, of course, on the profile shape.

In the following, the profiles of the Bragg reflections were fitted with Gaussian, Lorentzian, or pseudo-Voigt weighting functions. These shape functions can be differentiated by comparing the fits of the tails of the peaks. Nevertheless, they all lead to similar half-widths. The definitions of Dqi and DFi ensure that the same proportion of each reflection is included within the integration boundaries. Thereby, the inferred values from the fits are given with an uncertainty, including the small variation found when different models are used to describe the experimental profile shapes.

It has also been noted [START_REF] Millane | Relating reflection boundaries in X-ray fiber diffraction patterns to specimen morphology and their use for intensity measurement[END_REF] that eqns.7 and 8 are valid in the limit of small half-widths, which is the case here, as the analysis shows. On the other hand, the breadths of the reflections were found to be substantially larger than the beam size, as required by the method. Lastly, eqn.8 is valid for small b values, [START_REF] Millane | Relating reflection boundaries in X-ray fiber diffraction patterns to specimen morphology and their use for intensity measurement[END_REF][START_REF] Fraser | Digital processing of fibre diffraction patterns[END_REF] which is also found from the fits. Thus, by examining the radial and angular profiles of the reflections in the configuration where the incident beam is parallel to the microelectrodes (w = 0 in eqn.8), eqns.7 and 8 have been used to estimate the three morphological parameters of the thin film, namely llat, laxial, and b. The focus of the analysis was on some well-resolved reflections to avoid interferences by overlapping profiles of adjacent reflections. In addition, the reflections located at large wavevectors allowed the most reliable results to be obtained, as the diffracted path is less obscured by the microelectrodes. For illustration, the collected intensity around the spots (042 j ) and (132 j ) arises from more than 90% of the sample volume contained in the gaps between the electrodes (total scattering angle 2q > 17°).

For reflections with lower q values, the diffraction spots located along the meridian of the detector were used as the blind zone is minimized. Note also that our analysis was carried out with most of the spots located away from the meridian axis and at relatively high q. Indeed, for an incident angle equal to the critical angle (ai = ac) the patterns present a maximum in intensity localized at the exit angle ae = ac (Yoneda's peak). [START_REF] Yoneda | Anomalous surface reflection of X-rays[END_REF] This effect, well encountered upon reflectometry experiments, enhances the measured intensity close to meridian and to equator, principally at low q, but does not affect the other Bragg spots. So the potential residual contribution of the Yoneda's effect does not change our conclusions.

A selection of angular and radial profiles is given in Figure 5. Good fits as well as less satisfactory fits for some arc reflections are deliberately illustrated to give an idea of the uncertainty obtained in the results. Nevertheless, the power of our procedure, when a whole set of arc reflections is investigated for the same GIWAXS configuration, is that the analysis converges to inferred morphological parameters of the film with low uncertainty. To our knowledge this direct method outperforms all other means for investigating thin films in complex geometries, such as films composed of nanowires placed between microelectrodes of microelectronic circuits. Thus, the mean crystallite sizes and disorientation were determined to be llat = 140 ± 20 Å, laxial = 1140 ± 30 Å and b = 12.2 ± 0.5°. The low value found for the mean disorientation shows that the nanowires are aligned almost all parallel along the normal to the electrodes. The planar mosaïcity is ± 6° and the distribution of the nanowires confined to the 2D plane have an orientational order parameter of

𝑆 2D = 〈2cos * 𝛽 -1〉 = 0.91 ± 0.02
This value is in good agreement with the order parameter inferred previously from SEM images. [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] With S3D ≈ 0.986 (see above) and S2D ≈ 0.91 the nanowires do not only lie flat onto the substrate, but are also aligned parallel to the normal to the electrodes. This confirms the growth mechanism by supramolecular electropolymerization along the current lines of the electric field as reported previously. [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] Figure 5. Angular (a,b) and radial (c) profiles of some diffraction spots after subtraction of the scattered background intensity, for the configuration in which the incident x-ray beam is oriented parallel to the micro-electrodes, and a concentration of 2.5 mg/mL is used (see Fig. 2). The angular profiles (solid circles) can be fitted using (a) the sum of two Gaussian functions for the spot (042 j ) (full width at half maximum (FWHM) is 5.0 ± 0.3°) or (b) the sum of two Gaussian functions or a Lorentzian for the spot (110) with FWHM = 3.1 ± 0.2 °. (c) The radial profile (solid circles) of the diffraction spot (002) as a function of the wavevector q is composed of two peaks centered respectively at q = 1.2951 Å -1 and 1.3021 Å -1 . The presence of these overlapping peaks arises from the two preferential orientations of the nanowires on the substrate (in green and in blue; see Fig. 2), but also to the imperfections of the sample: depending on the position of the diffracting zone, either upstream or downstream, the apparent q values are slightly different. The profile can be fitted using the sum of two Gaussian functions of full width at half maximum (FWHM) = 0.0035 Å -1 and 0.0071 Å -1 respectively.

Crystallite sizes and coherence length. Analysis of very long anisotropic objects, such as fiber diffraction patterns, has pointed out that the "coherent molecular length" is a better concept than the axial crystallite size. [START_REF] Langford | Scherrer after sixty years: A survey and some new results in the determination of crystallite size[END_REF] Thus, the average length of the ordered crystallites along the long-axis of the nanowires is referred to as the coherence length. The intensity collected by spots, located along the equator of the detector, results from the layer planes of periodicity c in the real space and distant by l/(c sina) in the reciprocal space. Equation 7 can be rewritten by introducing the average coherence length of the molecules along the nanowire long axis, Lc, and using the Scherrer formula relating the breadth, or full width at half-maximum, of the diffraction spot (002), with respect to the wavevector to the average axial size of the crystallite along the nanowires:

𝐿 l = 2𝜋𝐾 Sch /Δ𝑞 CC* (9)
where the Scherrer constant, 𝐾 Sch , has a value of about 0.9. [START_REF] Langford | Scherrer after sixty years: A survey and some new results in the determination of crystallite size[END_REF][START_REF] Smilgies | Scherrer grain-size analysis adapted to grazing-incidence scattering with area detectors[END_REF] The intensity profile of the spot (002) along the equator of the detector (i.e. qx-axis of Fig. 2) is shown in Fig. 5c. By applying eqn.9 one finds Lc ≈ 1200 Å. This value is in good agreement with the axial crystallite size laxial = 1140 ± 30 Å inferred from the fits of the breadths of the other reflections (eqns.7 and 8). This remarkable length represents more than 200 molecular lengths along the long axis of the nanowires, grown by electropolymerization. Concerning the lateral crystallite size, laxial = 140 ± 20 Å, the inferred value from the fits appears larger than some narrow nanowires, that can be observed by SEM and AFM.

We note however that these techniques are more sensitive to a number-average rather than a weight-average as in x-ray techniques.

New information provided by GIWAXS on nanowire devices prepared by electropolymerisation. To sum up, the new type of GIWAXS analysis presented above is suitable to fully characterize TAA-based supramolecular nanowires directly electropolymerized between microelectrodes. The new information which is obtained compared to the more classical methods of analysis in the literature [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] is manifold and includes: i) the unequivocal determination of the full set of crystallographic cell parameters and the modelling of the molecular organization; ii) the specification of the orientation of the columnar structure with respect to surface and with respect to electrodes; iii) the quantification of the distribution of domain orientations with two order parameters; iv) the estimation of the correlation lengths along nanowires and in the cross-sectional plane. This rich information is obtained while the nanowires could only be probed in a single sample orientation due to the presence of electrodes obstructing the beam. All the parameters accessed here cannot be determined to date by other techniques, and they are all of crucial importance to further probe structure-function relationships in organic electronic devices.

CONCLUSION

In conclusion, we have demonstrated the power of the GIWAXS technique to fully characterize the structural and orientational parameters of supramolecular nanowires grown in situ in organic electronic devices. This was particularly challenging as the nanowires remain inserted into the microelectronic circuits. By taking TAA-based supramolecular polymers as an example, we have been able to confirm their molecular structure when trapped between interdigitated electrodes. In addition, we have determined their coherence length in such a configuration (> 100 nm). This is a quantitative result which remains inaccessible so far by other techniques. For instance, defects were sought by microscopy techniques such AFM. This was only possible on very large and hard monocrystals of triarylamines, and obviously only at their surfaces. Here, it was impossible to get this information by probing a very thin and soft nanowires configuration. Even if the fibers look perfect on large distances, information is lost on their internal structure. Therefore, it is a strong asset of the GIWAXS technique to access information that is not available in real space. Finally, both 3D and 2D orientation parameters have been determined, which is, once again, a clear advantage compared to microscopic analyses that are useful only for 2D parameters. With a remarkable precision, we demonstrate that the nanowires align not only parallel to the substrate, but equally along the current lines of the electric fields between the microelectrodes. As a whole, our results demonstrate that a single GIWAXS experiment can provide all the necessary quantitative information to characterize in situ the most important structural and orientational features that are needed to characterize supramolecular electronic devices. If this study also reinforces our previous claims on supramolecular electropolymerization, [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] it highlights more importantly the high potential of GIWAXS to become a reference method for achieving the full characterization of nanowiring processes involving various types of supramolecular polymers and polymer self-assemblies in complex geometries. With information that cannot be obtained by other means, our method could foster new supramolecular wiring approaches to be implemented, quantified, and validated in nanotechnologies.

EXPERIMENTAL SECTION

Materials. Synthesis and growth of triarylamine TAA 1 nanowires by electropolymerization were carried out following the procedures described previously. [START_REF] Ellis | Supramolecular Electropolymerization[END_REF][START_REF] Nyrkova | Supramolecular Self-Assembly and Radical Kinetics in Conducting Self-Replicating Nanowires[END_REF] A 1.25×10 7 V/m applied DC electric field was used, as it was shown to be optimal for the alignment of nanowires during the supramolecular polymerization of TAA 1. [START_REF] Ellis | Supramolecular Electropolymerization[END_REF] The samples were prepared on special glass substrates, composed of 15 interdigitated gold micro-electrodes pairs of 100 nm height, 2.5 mm long, having each a width of 10 µm and spaced away by 10 µm. The GIWAXS data were collected post-process.

In all the manuscript "in-situ" refers to the process of fabrication, meaning that nanowires were synthesized in the final device. The substrates with their microelectrodes were mounted on a precise goniometer platform allowing the alignment and movement of the sample according to six degrees of freedom with a very good repeatability (full vertical axis rotation: 1 mdeg). [START_REF] Fontaine | Soft interfaces and resonant investigation on undulator source: A surface X-ray scattering beamline to study organic molecular films at the SOLEIL synchrotron[END_REF] In the configuration where the incident beam is parallel to the microelectrodes (w = 0, Fig. 1), the misalignment (non-parallelism) of the incident x-ray beam with the 2.5 mm -long electrodes is estimated to be less than 0.3 degrees.

GIWAXS experiments.

The experiments were conducted at room temperature on the SIRIUS beamline at the French National Synchrotron Radiation facility SOLEIL in Saint-Aubin, France. [START_REF] Fontaine | Soft interfaces and resonant investigation on undulator source: A surface X-ray scattering beamline to study organic molecular films at the SOLEIL synchrotron[END_REF] An X-ray energy of 10 keV (X-ray wavelength 1.24 Å) was selected with a Si(111) double crystals monochromator. The beam size 500 × 100 µm 2 (H × V) was obtained using a two-slit system. A beam stop was used to stop direct and reflected beams. Two-dimensional x-ray scattering data were collected on a DECTRIS PILATUS3 1M 2D detector positioned 472 mm downstream of the sample. The diffraction pattern of silver behenate was used to calibrate the sample-to-detector distance. [START_REF] Nyam-Osor | Silver behenate and silver stearate powders for calibration of SAS instruments[END_REF] Two laterally rotated detector positions were used for covering the scattering vector range up to 2 Å -1 . Patterns were corrected from silica glass and air scattering by subtracting normalized background obtained from empty cell. For the data presented, the incidence angle was set at 0.12°, below the total external reflection critical angle of the glass substrate. Typical exposure times for these samples ranged from 1 to 20 min. Because qz is not probed for more than the very shallow exit angles, the illuminated detector pixels do not linearly correspond to the examined qspace. Therefore a new mesh was generated and processed starting from the raw image of the 2D detector to lead to the q-map shown in the GIWAXS patterns (Fig. 2, Fig. S1). The intensity I(q)

profiles were obtained from background subtracted images using home-developed software, through re-sampling to polar I(q, f) coordinates and renormalization to entire integration range going from f = 0° to f = 90°.

Measurement of Bragg reflection intensities.

Accurate subtraction of background is important to obtain accurate intensity measurements. A first approach used a global background function expanded as a polynomial. However, the scattered background is not circularly symmetric. Thus, to obtain reliable radial and azimuthal profiles of the Bragg reflections, two-dimensional background level matrices were obtained from peripheral regions of reflection arcs using series of one-dimensional ring profiles. This was particularly advantageous for estimating the radial profiles when there is little space between close arc-shaped reflections of adjacent Bragg spots.

For an ensemble of nanowires exhibiting preferred orientation, the scattered intensities are modeled as an incoherent sum of intensities from the individual objects. The intensity at a given point in reciprocal space having spherical coordinates (qi,yi,si) results from the product of weighting functions accounting for radial peak broadening (in q, distribution function Gq centered at the maximum qi with width Dqi) and preferred orientation (functions GY and Gs centered at the maxima yi and si and having widths Dyi and Dsi respectively). Summing the weighted contributions to Ii from all reciprocal lattice points n (provided that they intersect the Ewald sphere), the expression for the intensity reads 𝐼 S (𝑞 S , 𝜓 S , 𝜎 S ) ∝ p 𝐶 corr 𝐺 Z E r𝑞 S -𝑞 E , Δ Z s t 𝐺 u r𝜓 S -𝜓 E , Δ u s t 𝐺 v r𝜎 S -𝜎 E , Δ v s t (10) where Ccorr contains various geometric corrections in the GIWAXS pattern obtained with a sample tilted by an angle q0 (in the present experiments, q0 = 0.12°) with respect to the incoming beam (see Fig. 1e): The first factor is the polarization factor [START_REF] Smilgies | Geometry-independent intensity correction factors for grazingincidence diffraction[END_REF][START_REF] Schlepütz | Improved data acquisition in grazing-incidence X-ray scattering experiments using a pixel detector[END_REF] at the linearly polarized SIRIUS beamline at the SOLEIL synchrotron (the fraction of radiation that is polarized in the horizontal direction, pH, is larger than 98%), and the second factor is the absorbance correction. Here, 2q designates the total scattering angle, while the angles µ and d define the direction of the diffracted beam with respect to the incident beam: µ is the horizontal ("longitude") angle and d the vertical ("latitude") angle of the intersection of the diffracted beam with the Ewald sphere of reflection. In this experimental work, the detector is vertical, and thereby the calculation of µ depends on the angular position of the detector, while d does not. The absorbance correction results from the longer path of x-rays at low exit angles than at high angles, through a film of thickness T, assumed here to have a uniform absorbance. Nevertheless, as the sample film is very thin (T~100 nm), with an extinction coefficient µa ~ 2 mm -1 for a wavelength of 1.24 Å (calculated from the chemical composition and tabulated atomic values), this absorption correction is negligibly small.

𝐶 corr =

Figure 1 .

 1 Figure 1. (a) Molecular structure of TAA 1; (b) Generic mechanism of supramolecular polymerization of TAA 1 through an oxidation / nucleation / growth process; (c and d) Atomic Force Microscopy (AFM) images of nanowires [reproduced from ref.26 with permission] formed by supramolecular electropolymerization of TAA 1 solutions in heptane at 2 mg/mL (c) and 5 mg/mL (d); the solutions were drop casted onto 2.5 mm long interdigitated gold micro-electrodes (width 10 µm; height 100 nm; spaced away by 10 µm) at a 1.25×10 7 V/m applied DC electric field 26 ; (e) Schematic of the GIWAXS setup on the SIRIUS beamline at the French National Synchrotron Radiation facility SOLEIL (Saint Aubin, France)[START_REF] Fontaine | Soft interfaces and resonant investigation on undulator source: A surface X-ray scattering beamline to study organic molecular films at the SOLEIL synchrotron[END_REF] including relevant geometry definitions. The sample is tilted by 0.12° with respect to the incident beam for grazing incidence wide-angle measurements. The sample can be rotated about its vertical axis: its position is defined

7 ( 3 )

 73 Å, b = 21.0(6) Å and c = 9.7(0) Å, and angles a = 94.(5)°, b = g = 90°.

Figure 2 .

 2 Figure 2. (a) GIWAXS pattern with the incident x-ray beam aligned parallel to the electrodes (w=0; Fig.1e), for an electropolymerized 2.5 mg/mL TAA 1 solution on interdigitated gold microelectrodes. The indexation of diffraction spots (hkl indices on the right of the spots) indicates two main preferential orientations (in green and in blue); (b) the corresponding orientations of reciprocal cell axes. The c*-axis is oriented parallel to the substrate while the a*-axis is either tilted by 0-30° from (green) or normal to (blue) the substrate; (c) The indexed radial profile leads to a monoclinic cell with a=33.7(3) Å; b=21.0(6) Å; c=9.7(0) Å; a=94.(5)°; b=g=90°.

Figure 3 .

 3 Figure 3. Schematic representations of the molecular organization. Left: association of triarylamine segments by H bonding and cross-section of double-strands; bottom left: structure of triarylamine double-strands from side and equivalent columns of piled molecular units, showing periodicity h of two molecular thicknesses hmol along strand axis*. Right: Arrangement in the columnar plane according to the lattice geometry and systematic extinctions; bottom right: Longitudinal positioning of columns and monoclinic structure. The aliphatic chains fill the empty spaces between the columns.

  2): the rectangular sublattice a x b, with a non-orthogonal angle a between the b-axis and c-axis (a = 94.5°), is oriented with either one direction tilted, or normal to the substrate. The relative intensities of the observed spots indicate that the distribution of the crystallites with the planes normal to the [110] direction of the monoclinic cell (spots in green in Fig.2a) is the preferred orientation, compared to the second orientation with reticular planes normal to the [100] direction (spots in blue). The two preferential orientations revealed at low concentration underline that the alignment allows the columns composing the crystals (Fig.3) to build the set of the more dense reticular planes along the glass substrate. These planes are those where the columns are the closest to each other. The reticular planes that are the most dense, are normal to the diagonal [110] of the rectangular sublattice a x b.

  -30cos * 𝜙 + 3〉 = 0.951 ± 0.006.

  positioned and randomly rotated about the c-axis. In such a case, the recorded diffraction pattern is that diffracted by a single crystallite object, cylindrically averaged about the Z axis in reciprocal space. As a result, the positions of the Bragg reflections are described by the cylindrically projected reciprocal lattice onto the plane of the 2D-detector, provided the Ewald condition of diffraction is fulfilled. The cylindrical polar coordinates of the Bragg reflection in the reciprocal space are noted (qi,si) with

  𝐾 𝜆 𝐷 𝑢|sin 𝜎 S | + 𝑣|cos 𝜎 S | (|sin 𝜎 S | + |cos 𝜎 S |)|cos 𝜃 S | cos * 2𝜃 S (7) and Δ𝜙 S = 𝛽 + 2𝜋 𝑢|cos 𝜎 S | + 𝑣|sin 𝜎 S | 𝑞 S (|sin 𝜎 S | + |cos 𝜎 S |) |sin 𝜎 S | g(cos 𝜃 S cos 𝜔) * -(cos 𝜎 S -sin 𝜃 S sin 𝜔) * (8)

  Additional GIWAXS patterns at high concentration are presented. Corresponding Authors: Nicolas Giuseppone -Institut Charles Sadron, Université de Strasbourg, C.N.R.S., UPR22, 23 rue du Loess, 67034 Strasbourg, France; Email: giuseppone@unistra.fr ; ORCID: 0000-0003-4093-3000 Patrick Kékicheff -Institut Charles Sadron, Université de Strasbourg, C.N.R.S., UPR22, 23 rue du Loess, 67034 Strasbourg, France; Synchrotron SOLEIL, Saint Aubin, L'Orme des Merisiers, BP48, 91192 Gif-sur-Yvette, France; Email: patrick.kekicheff@ics-cnrs.unistra.fr ; ORCID: 0000-0001-8465-9627 Authors Adress: Melodie Galerne -Institut Charles Sadron, Université de Strasbourg, C.N.R.S., UPR22, 23 rue du Loess, 67034 Strasbourg, France ; ORCID: 0000-0002-8535-2923 Benoît Heinrich -Institut de Physique et Chimie des Matériaux de Strasbourg, Université de Strasbourg, C.N.R.S., UMR7504, 23 rue du Loess, 67034 Strasbourg, France ; ORCID: 0000-0001-6795-2733 Arnaud Hemmerle -Synchrotron SOLEIL, Saint Aubin, L'Orme des Merisiers, BP48, 91192 Gifsur-Yvette, France ; ORCID : 0000-003-1722-2024 Philippe Fontaine -Synchrotron SOLEIL, Saint Aubin, L'Orme des Merisiers, BP48, 91192 Gifsur-Yvette, France ; ORCID : 0000-0003-3394-6508

  w𝑝 H y1sin * 𝜇 cos

	*	𝛿| + r1 -𝑝 H sin * 𝛿t}	1 -𝑒 𝜇 ‡ 𝑇 y •€ • '& sinƒ " 4 1 sin𝜃 C + sin𝛿 | … 4 ) sin † 1	(11)

ACKNOWLEDGMENT

We thank SOLEIL synchrotron for beamtime provision. Experiments were carried out within the framework of the MICASOL platform set between Institut Charles Sadron, C.N.R.S., Strasbourg, France, and SOLEIL synchrotron (www.ics-cnrs.unistra.fr/micasol).

Part of the experimental work (sample preparation) was performed within the framework of the IRTG "Soft Matter Science: Concepts for the Design of Functional Materials" (M.G. and N.G.).

AUTHOR INFORMATION

Notes

The authors declare no competing financial interest.