C Chassagneux

M Descorme

L Besson

Khrouz

F Schmit

L Bois

R Chiriac

F Toche

F Chassagneux

C Descorme

M Besson

L Khrouz

Porous microspheres of manganese-cerium mixed oxides by a polyvinylpyrrolidone assisted solvothermal method

Keywords: Crystallites, Segregation, Hydrothermal crystal growth, Nanomaterials, Oxides

Porous microspheres of manganese-cerium mixed oxides by a polyvinylpyrrolidone assisted solvothermal method

Introduction

MnO x -CeO 2 materials have been reported as highly active catalysts in many different oxidation reactions [1]. It is thought important to develop some new methods to synthesize MnO x -CeO 2 materials with a controlled morphology. In cerium oxide, cerium may easily switch between the III and IV oxidation states and might be substituted with lower valence elements, inducing the creation of oxygen vacancies. Cerium oxide is also known to release oxygen at high temperature. CeO 2 is used as fast ion conductor, oxygen storage capacitor, catalyst and electrolyte for solid oxide fuel cells [2][3].

Nano-structured ceria synthesis has been recently reviewed [3,4]. By changing the temperature and the NaOH concentration or by using organic additives which ones adsorb onto the surface of the growing particles, different morphologies of ceria could be obtained [5][6][7][8]. For example, the use of polyvinylpyrrolidone (PVP) as a capping ligand has been reported to allow the formation of monodisperse spherical ceria particles with size in the range 100-500 nm [9][10][11][12][13][14][15][16][17][18][19]. Manganese oxides with different Mn valencies, are also well known in oxidation catalysis and as natural oxidants [20]. A lot of research is devoted to the elaboration of nano-structured manganese oxides-based catalysts for the oxidation of volatile organic compounds [21], the selective reduction of NO with NH 3 [22] or the oxidation of water [23]. The combination between manganese and cerium oxides has been the subject of several studies, mostly in oxidation catalysis [1,[24][25][26][27][28][29][30][31][32], but also for the selective reduction of NO by NH 3 [33,34]. A synergetic effect is often reported between manganese and cerium. Manganese is known to easily substitute cerium in the ceria lattice, increasing the oxygen mobility in ceria. The enhancement of the catalyst efficiency is then explained by the improvement of the oxygen activation and transfer through redox cycles involving both manganese and cerium. The reduced manganese oxide is re-oxidized by ceria and the reduced cerium oxide is re-oxidized to CeO 2 by the oxygen from the air [27]. Structural characterization studies have been realized [35][36][37]. Recently, theoretical studies are gain more and more interest to explain the behaviour of such mixed oxides [38][39][40][41]. Ce-Mn oxides are also used to benefit from their specific electrical properties, such as in solid oxide fuel cells [42][43][44]. Ce-Mn oxides can be synthesized by impregnation, coprecipitation, sol-gel process, hydrothermal, or glycothermal and combustion methods [27]. Usually, a complex mixture is formed containing both a solid solution and separated manganese oxides phases. Only a few studies concerned the synthesis of mixed manganese-cerium systems via solvothermal methods, generally allowing a better morphology control, as cubes, rods or flowers, using either pH control or acrylic acid adding [24,25]. In this paper, we describe the synthesis of porous microspheres of manganese-cerium mixed oxides using a solvothermal method and PVP, as a capping ligand. Three different Mn:Ce ratios were considered with an increasing amount of PVP. The final calcined products were studied by X-ray diffraction, SEM, TEM, EPR, Raman spectroscopy, TGA and nitrogen adsorption desorption.

Experimental Section 1.1 Synthesis of the materials.

Ce(NO 3 ) 3 .6H 2 O (5 mmol, 2,17 g) was mixed in ethanol (30 mL) and water (10 mL) with 1 g (9 mmol) or 3 g (27 mmol) of PVP (MW 40000). The manganese salt Mn(NO 3 ) 2, .4H 2 O was added in various proportion. The Mn:Ce ratio was fixed at 0, 0.25, 0.5, or 1. The solvothermal treatment was done in a 250 mL Teflon-lined autoclave, at 200°C (10°C/min), for 48 h. The solid was separated by filtration and washed three times with water, before to be dried at 70°C. A brown powder was obtained whenever manganese nitrate was used. Finally, a thermal treatment under air was realized at 400°C at 1°C/min during 5 h to eliminate the polymer.

The hydrothermal stability was tested upon treatment at 200°C during 24 h. 80 mg of sample was introduced in a 250 mL Teflon-lined autoclave containing 100 mL of water. Black powders were obtained after the hydrothermal treatment. Samples were named as CM x P y where x = Mn:Ce ratio and y = polymer weight (y= 0, 1 or 3 g). Finally, if the sample was hydrothermally treated, an h was introduced at the end of the sample name (CM x P y -h).

Characterization methods

X-ray diffraction (XRD) patterns were recorded on a PANalytical Xpert Pro diffractometer (The Netherlands) equipped with a monochromator, using Cu Kα radiation (wavelength of 1.54 Å). The width of the (220) reflection of ceria was used to calculate the crystallite size because there was no oxide manganese reflexion in this zone. SEM pictures were acquired on a S800 Hitachi microscope operating at 15 kV in the "Centre Technologiques des Microstructures de l'Université de Lyon", after gold/palladium metallization. TEM pictures were acquired on a Topcon instrument operating at 200 kV in the "Institut Lumière Matière, University Lyon 1". Prior observations, the samples were ground and deposited directly on a copper grid coated with a holey carbon film. EPR measurements on selected samples were carried out using a X-band spectrometer (Bruker EMX plus, double mode) and a X-band (9.4 GHz) Bruker ESP 500E spectrometer using a standard rectangular cavity. Simulations of the spectra were realized using easyspin (Matlab toolbox). The experimental spectra were simulated using the following parameters: frequency= 9.389 GHz; temperature= 25 K; microwave power= 21 mW; amplitude modulation= 3 G; and modulation frequency= 100.00 kHz. Raman spectra were acquired on a Renishaw RM 1000 microspectrometer at λ=532 nm (CECOMO, University Lyon I).

Thermogravimetric analyses of the calcined samples were realized with a TGA/SDTA 851 Mettler Toledo apparatus. 20 mg of sample were placed in the platinum crucible. Air was used as the reactive gas (50 mL. min -1

). All samples were heated at 20°C.min -1 from 30 to 1200°C. First derivative curve was used to calculate the different mass loss steps, while the SDTA curve (differential thermal analysis) provided information about the thermal phenomena occurring upon the sample heating.

Textural characterizations were realized using nitrogen adsorption/desorption isotherms on a BelsorpMini (Bel Instruments, Japan). Prior analysis, samples were outgassed under vacuum (100°C for 4 h). The pore size distribution (max radius) and the mesoporous volume were calculated from the desorption branch of the isotherms using the Barrett-Joyner-Halenda (BJH) method. The total porous volume was measured at P/P 0 =0.98. The surface area was determined by the Brunauer-Emmett-Teller (BET) approach.

Results

After the solvothermal treatment.

At the end of solvothermal treatment, the polymer capping the inorganic phase is evidenced by TEM (Fig. 1a &b). The departure of the polymer was followed using TGA/SDTA analysis (Fig. 2). A weight loss of 13.1 % was measured at 250°C (Fig. 2a), associated to an exothermic peak (Fig. 2b), due to the polymer oxidation process. The weight loss measured, has allowed an estimation of the amount of PVP around the oxide at the end of the solvothermal process. It is found the 0.25 PVP unit for one cerium was present. All the following analysis have been done on samples thermally treated at 400°C, in which the polymer has been eliminated.

XRD results

The X-ray diffraction patterns of the CM x P y samples are given in Figure 3. The influence of the Mn:Ce ratio on the ceria structure is shown in Figure 3a. In fact, a typical well-crystallized ceria pattern with a cubic fluorite structure was obtained when no manganese was added (CM 0 P 3 ). The width of the half maximum height of the (220) reflection at 47.5° was equal to 0.58°, corresponding to a crystallite size of 16 nm. Upon addition of manganese, CM 0.25 P 3 and CM 0.5 P 3 patterns revealed a less crystallized ceria, with a shift of the (220) reflection to higher angles, from 47.5° to 47.8°. The shift of the (111) reflection is also seen in the insert Fig. 3a. The width of the (220) reflection, 1.61° and 2.24°, respectively indicated a 5.8 and 3.9 nm crystallite size. The hausmannite phase Mn 3 O 4 could be detected in the CM 0.5 P 3 sample from the (202) reflection at 36.2° (ICDD n° 89-4837). For the highest Mn:Ce ratio (1), the ceria phase looked the same as in the CM 0.5 P 3 sample, while the hausmannite reflections became more intense. The width of the (220) ceria reflection was found constant (2.20°). In the reference sample, MP 3 , prepared without cerium, a well-crystallized hausmannite phase was obtained. The shift of the ceria reflection revealed that part of the manganese ions was inserted inside the ceria structure. Such decrease of the cell parameter of ceria upon insertion of the manganese has already been reported [45,47]. In fact, a Rietveld analysis has recently confirmed that when Mn 2+ (ionic radius 0.67 Å) substitutes cerium in ceria (Ce 4+ , ionic radius 0.92 Å), a decrease of the cell parameter occurs, as well as a decrease of the crystallite size [47]. The solubility of Mn 2+ inside the ceria lattice (at 1300°C) has been estimated between 5 and 10 mol % [44]. The smaller Mn 4+ cation (Mn 4+ , ionic radius 0.53 Å) is also expected to be easily inserted inside the ceria [48]. Figure 3b shows the influence of the proportion of polymer on the ceria structure for the CM 0.5 P x samples. The crystallinity of the ceria (crystallite size=6.8 nm) and hausmannite phases were found to be a little higher in the CM 0.5 P 0 pattern compared to the CM 0.5 P 1 (crystallite size=5.5 nm) and CM 0.5 P 3 samples (crystallite size=3.9 nm). Moreover, in the CM 0.5 P 3 sample, a slight shift of the (111) reflection towards to higher angles was observed, indicating that the degree of substitution of manganese inside ceria was higher at higher polymer ratio. By increasing the amount of polymer, the substitution of manganese into the ceria lattice (i.e. solid solution) was favoured and the Mn 3 O 4 crystallization was prevented. Indeed PVP acts as ligand and form complexes with metal ions through its amide groups. PVP covers the small nanocrystals ceria seeds and limit the crystal growth [2]. Manganese ions diffusion inside the ceria seeds is favored by the presence of PVP. This explains why manganese-cerium solid solution formation is favored and crystallization prevented.

SEM and TEM results

The amount of polymer also influenced the crystallites morphology. The SEM pictures for CM 0.5 P 0 (Fig. 4a) and CM 0.5 P 1 (Fig. 4b) revealed the important agglomeration of micrometric grains, while large micrometric spheres with a hole were only obtained in the case of the CM 0.5 P 3 sample (Fig. 4c). The Mn:Ce ratio also strongly influenced the morphology. In fact, submicronic spheres were observed in the samples with the lower Mn:Ce ratios: 0 (Fig. 4d) or 0.25 (Fig. 4e); while large micrometric spheres with a hole, were obtained only in the CM 0.5 P 3 sample (Fig. 4c). Moreover, in the sample, obtained without any cerium (MP 3 ), the hausmannite Mn 3 O 4 phase exhibited a flower-like morphology of about 2 microns in size (Fig. 4f). Submicronic spherical morphologies of ceria were already noted in the assisted PVP hydrothermal synthesis of ceria [14,17,18,49,51]. Indeed, PVP seems to easily deposit on the surface on the ceria seeds and stabilize the nanosphere morphology [2].

TEM pictures showed that in the absence of manganese (CM 0 P 3 sample), regular 300 nm spheres were obtained (Fig. 5a). For a Mn:Ce ratio of 0.25, similar spherical aggregates were observed (Fig. 5b), but with a lower size in the range 100-200 nm. At higher manganese ratio, (CM 0.5 P 3 sample) (Fig. 5c), the spherical aggregates usually showed a diameter below 100 nm and looked like snowflakes, exhibiting a porous structure. At a larger scale, these 100 nm spherical type particles were aggregated into large micrometric spheres, as observed by SEM. For a Mn:Ce ratio of 1, snowflake's like spheres were also observed (Fig. 5d). Without PVP (Fig. 5e), the snowflake morphology is not found as seen in the picture of the CM 0.5 P 0 sample. Using a lower amount of PVP (sample CM 0.5 P 1 shown in Fig. 5f), a snowflake morphology is obtained, but their size is less well defined.

The HRTEM images revealed that the ceria spheres in CM 0 P 3 (Fig. 6b) consisted of aggregates of smaller crystallites of ceria, of rectangular ones (15 x 30 nm). For a Mn:Ce ratio of 0.25 (CM 0.25 P 3 ), the spherical aggregates (Fig. 6e) was formed with even smaller crystallites (7 nm). In the presence of a higher manganese ratio (CM 0.5 P 3 sample) (Fig. 6h), the aggregates were constituted of the smallest ceria crystallites (3 nm). Electron diffraction patterns revealed a very clear preferential orientation of the ceria crystallites, in the CM 0 P 3 and CM 0.25 P 3 samples (Fig. 6c and6f). The preferential orientation of each ceria aggregate in these two samples was due to the stacking of the small crystallites with the same orientation. In fact, electron diffraction patterns obtained from a single sphere consisted of well-defined spots, like a single-crystal pattern, with a definite zone axis. By tilting the sample in a large angular range (+-60°), different zone axis could be selected, demonstrating the ordered rigorous stacking up of the crystallites, which was not of spherical symmetry. Crystallites are stacked parallel to the {111}, {110} and {100} planes with the lower surface energy [51][52]. In the CM 0.5 P 3 sample, crystallites started to lose their preferential orientation, as shown by the presence of bows and the formation of rings in the electron diffraction pattern, although obtained from a single sphere (Fig. 6a).

EPR characterization

EPR spectra, with a signal at g=2 and a hyperfine structure were characteristic of Mn II [53] (Fig. 7a). Both the absence of any change with temperature (blue curve at 15 K and green curve at 25 K) and the presence of the hyperfine structure indicated that Mn II species were isolated and without any strong interaction with the other Mn II species. Moreover, at least two types of Mn II were detected, revealing two different distances between two Mn II centers. The coordination environment of Mn II is a distorted octahedral. The presence of Mn II species was explained by the substitution of cerium with manganese in ceria. The Mn II species were found to be more present in the CM 0.5 P 3 sample (blue line) than in the CM 0.5 P 0 sample (red line). The EPR analysis confirmed the XRD observations relative to the easier formation of the solid solution in the presence of the polymer.

Further EPR experiments realized in perpendicular mode (Fig. 7b) confirmed the presence of at least two types of isolated Mn II species located inside ceria, with low interaction with other Mn II in the ceria framework. A signal characteristic of Mn IV at g=4, with a lower resolution, was also observed [54].

Moreover, additional information was obtained from the simulation of an experimental spectrum obtained in parallel mode (Fig. 7c). Simulations of the spectrum were done with the following parameters: S =1/2, 65%, g = [2. 16 , E/D=0.05 for Mn IV species, where a corresponds to the hyperfine interaction tensor (electron-nucleus) and D and E are the zero field interaction tensors (high spin, E=0, axial, E#0, rhombic asymmetry). The presence of Mn III species could not be excluded. Indeed, Mn III is difficult to observe by EPR in oxides. So, Mn 3 O 4 , as attested by the XRD data could not be observed by EPR as shown by Lopez-Ortega & al. [55]. In addition, no trace of superoxide radicals, which are usually characterized by a sharp signal at g=2, was evidenced. Finally, these simulations allowed to estimate the ratio of the Mn II to Mn IV signals to 1.85 in the CM 0.5 P 3 sample. Mn IV represents 35% of the EPR signal, while Mn II represents 65% of the EPR signal. A broad signal at g=2.6 was also observed in the blue spectrum acquired at 25K (blue line, Fig. 7c). The simulation performed with only one undistorted Mn II species (D =0) could not reproduce this broad signal. Two Mn II species had to be considered: one undistorted Mn II (D =0) and another type of Mn II with some distortion (D#0).

Raman analysis

Raman spectra are presented in Figure 8. Only one peak was noticed at 463 cm -1 in the pure ceria samples (CM 0 P 0 and CM 0 P 3 samples Fig. 8a &8b). This peak was attributed to the symmetric breathing mode of oxygen atoms around the cerium ions in the cerium oxide cubic-phase [45]. The presence of the polymer resulted in an increase of the FWHM of the Raman bands, associated to a decrease of the crystallite size. Without PVP and upon manganese addition, for a Mn:Ce ratio of 0.5 (Fig. 8c), a broad vibration was noted at 463 cm -1 with a second one at 620 cm -1

. This second one vibration is related to the stretching mode of manganese-oxygen bound. For a Mn:Ce ratio of 1 (Fig. 8e), two vibrations were seen at 463 cm -1 and 650 cm -1

. In the case of PVP addition and upon manganese addition, for a Mn:Ce ratio of 0.5 (Fig. 8d), a broader vibration band was observed at 450 cm -1 with a second one at 630 cm -1

. While for a Mn:Ce ratio of 1 (Fig. 8f), two vibrations are noted at 450 cm -1 and 645 cm -1

. Then, in the case of PVP addition, the shift towards lower wavenumbers of the ceria vibration from 463 to 450 cm -1 might be related to the formation of a solid solution [47]. As it has been already noted from XRD analysis, PVP plays a decisive role in the formation of the solid solution between manganese and cerium. Upon adding manganese, for a Mn:Ce ratio of 0.5, an increase of the full width at half maximum (FWHM) of the ceria vibration band (at 450 cm -1 or 463 cm -1

) was also observed. Similar observations have been made in the case of Cu-Ce-O materials: an increase of the FWHM of the Raman bands was observed when the copper content increased. Such increase has been related to an increase in the concentration of the oxygen vacancies and to a decrease of the crystallite size [56]. The second vibration at 640 cm -1 was due to the Mn-O stretching mode in manganese oxides, revealing that part of the manganese species was not inserted inside ceria [47].

However, it was difficult to distinguish between MnO 2 and Mn 3 O 4 . In fact, vibration bands characteristic of MnO 2 would appear at 602 and 648 cm -1 while the one related to the Mn 3 O 4 is expected at 654 cm -1 [47]. We can also observe that the intensity of the Mn-O stretching mode is higher when no PVP has been added which may also been explained by the fact that the polymer has prevented the manganese oxide crystallization and has favored the solid solution formation.

Thermal analyses

TGA/SDTA experiments were realized under air (Fig. 9) The TGA curve for MnO 200°C till 1200°C, similar to the TGA profile obtained in the case of pure ceria [57]. TGA/SDTA of CM 0.25 P 3 showed the same progressive weight loss (Fig. 9c) as well as very small losses around 500 and 900°C. One possible interpretation is that most of the manganese was present as Mn II, probably because manganese was essentially inside the ceria lattice, but some Mn IV was also present in the sample as attested by the EPR analysis. TGA/SDTA of CM 0.5 P 3 showed a first loss below 200°C due to the loss of water, a second marked weight loss at 450°C, and a third marked weight loss at 900°C (Fig. 9c). By analogy with the behaviour of the MnO 2 sample (fig. 9a), the second loss was attributed to the reduction of Mn IV into Mn III, and to the formation of Mn 2 O 3 . Moreover, the third weight loss at 900°C was associated to the reduction of Mn III species into Mn II and the formation of Mn 3 O 4 . This analysis seems to indicate that, part of the manganese was involved in a solid solution with ceria while another part was located outside of the ceria particles, since two distinct reduction steps of the Mn IV species into Mn III and Mn II have been observed, associated to two endothermic signals. This analysis also suggested the presence of Mn IV species in the CM 0.5 P 3 , as attested already by the EPR analysis. These Mn IV species might be inserted inside ceria or located at the ceria surface since no MnO 2 phase could be detected by XRD or TEM analyses. The 0.4% weight loss at 900°C, measured in the case of the CM 0.5 P 3 sample, attributed to the reduction of Mn 2 O 3 into Mn 3 O 4 , was used to estimate the fraction of Mn IV in the sample, which might be considered as "free" (not inserted inside the ceria lattice). The following composition could be estimated for the fresh sample: Ce Mn(reduced) 0.2 O 2.2 + 0.15 Mn 2 O 3 . After treatment at 900°C, the composition would be Ce-Mn(reduced) 0.2 -O 2.2 + 0.10 Mn 3 O 4 . It is suggested that the initial composition is not far from Ce Mn(reduced) 0.2 Mn(oxidized) 0.3 O 2.8 , in which the reduced of Mn can be either inside ceria or as free Mn as Mn 3 O 4 . The TGA analysis of the CM 1 P 3 sample (Fig. 9c) was found to be somehow an intermediate between the CM 0.5 P 3 and the CM 0.25 P 3 samples. This result can be understood considering that fraction of manganese entered inside ceria and that the manganese in excess is mostly present as the hausmannite Mn 3 O 4 phase, indicating that the overall manganese oxidation state was lower. The 0.3% weight loss at 900°C was coherent with a higher proportion of Mn in a reduced state, compared to the CM 0.5 P 3 sample. This results showed that the Mn:Ce ratio had a strong influence on the overall manganese oxidation state. The presence of cerium IV certainly contributed to the stabilization of manganese IV. At low Mn:Ce ratio, Mn was found either as Mn II in ceria or as Mn IV. At intermediate Mn:Ce ratio, manganese was still stabilized as Mn IV, but Mn IV could be reduced at high temperature. However, when the Mn:Ce ratio further increased, manganese was mostly stabilized in a reduced state (hausmannite Mn 3 O 4 ).

Textural analyses

Textural characterizations were carried out using nitrogen adsorption/desorption isotherms.

The results are summarized in Table 1 and Figure 10. Specific surface areas were usually around 60 m². g -1 except for the CM 0.5 P 3 and CM 0.5 P 1 samples where higher specific surface area around 100 m². g -1 was obtained (Fig. 10a &10d). The hysteresis loop indicated the presence of mesopores.

Using the Barret-Joyner-Halenda (BJH) model, the pore size distribution could be derived from the desorption branch. The maximum pore radius was around 11 nm and the mesoporous volume was equal to 0.21 cm

3

. g -1

. The higher porosity observed for the CM 0.5 P 3 and CM 0.5 P 1 samples was in agreement with the snowflake morphology observed by TEM (Fig. 5c &5f).

Hydrothermal treatment

Such manganese-cerium mixed oxides could potentially be used as catalyst for the treatment of aqueous effluents in a process which is called catalytic wet air oxidation (CWAO) [28]. In such process, pollutants in water are eliminated using oxygen at high pressure and temperature. Before to study such catalytic process, we need to know if these compounds are stable under hydrothermal conditions. The hydrothermal stability of the CM 0.5 P 3 compound treated at 400°C was evaluated upon heating the sample at 200°C in water for 24 h. The sample changed from brown to black. X-ray diffraction pattern (Fig. 11a) revealed the formation of a crystallized bixbyite phase (Mn 2 O 3 ) characterized by the apparition of the (222) reflection at 32.9° (ICDD n° 01-089-4836). After hydrothermal treatment, the solid solution MnCeO crystallinity is improved since XRD intensity is higher. The influence of this hydrothermal treatment on the crystallite size (peak width) and on the textural properties is slight. Indeed the textural properties are preserved (Fig. 10b and Table 1) even if the mesoporous volume and the pore size have been a little increased. The micron-sized spherical morphology with a hole is conserved but whiskers of Mn 2 O 3 were observed by SEM on the surface of the CM 0.5 P 3 -h sample (Fig. 11b). The crystallite assembly always lead to snowflakes of 100 to 200 nm diameter (Fig. 11c &11d), which aggregate in spheres with a hole (Fig. 11b), as they were before the thermal treatment.

Discussion

Upon solvothermal treatment of a mixture of cerium and manganese nitrates, in the presence of polyvinylpyrrolidone (PVP), different solids with different morphologies could be prepared. The final product characteristics were shown to be highly dependent on the reaction conditions (Mn:Ce ratio and PVP concentration). Adding the PVP polymer which is a well-known capping agent for ceria seeds has a decisive influence on the morphology and structure of the compounds obtained. XRD and Raman analysis have shown that it has favoured the formation of a solid solution of manganese and cerium, probably by promoting the diffusion of manganese around the ceria seeds. In the absence of manganese, i.e. in the case of pure ceria, "spherical-type" objects (300 nm in diameter) were formed, with a very high textural organization. The organization of the crystallites induced by the presence of PVP has already been noticed as well as the formation of "spherical-type" objects [58]. A remarkable orientation of the crystallites occurs, leading to the construction of larger size (300 nm) spherical edifices. At first, the Oswald ripening mechanism, dissolution of the smaller particles and the recrystallization of larger particles, has been proposed to explain the formation of such morphologies [14]. However, a spontaneous self-organization of adjacent particles, inducing the sharing a common crystallographic orientation, might also be involved [10]. This mechanism called the "oriented attachment" has recently been shown to be predominant [9,10,18,19]. The morphology observed in our study looks very similar to what has been obtained by Zhou et al., upon synthesis in water at 140°C [18]. These authors obtained submicronic spheres formed by the aggregation of 5 nm crystallites. They have also reported a correlation between the sphere's size and the PVP/Ce ratio or the cerium concentration. On the opposite, our result differed from what was obtained by Chen et al., whom noticed the formation of 100 nm octahedral nanocrystals, upon synthesis in a mixture of water and ethanol at 160°C [14]. Such difference might be explained by the higher cerium and PVP concentrations involved in our study.

At low manganese concentration, Mn:Ce=0.25, a solid solution Mn-Ce oxide was obtained with "spherical-type" objects of 100 nm in diameter. These spheres consisted of smaller ceria crystallites, still remarkably ordered. EPR characterizations showed that (i) isolated Mn II was found inside the ceria lattice and that (ii) Mn IV was also present, localized either inside or near the ceria surface. TGA analysis revealed that such Mn IV species were not easily reducible. For a Mn:Ce ratio equal to 0.5, a MnCeO solid solution was mixed with a reducible Mn IV containing phase, resulting either from the Mn IV insertion inside ceria or from the "adsorption" of Mn IV species at the ceria surface. This later suggestion derives from the EPR analysis indicating the presence of Mn IV and the TGA analysis demonstrating that these Mn IV species were as reducible as in "free" MnO 2 . At such high manganese ratio, the ceria crystallite size decreased (3 nm) and the ordering of the crystallites also decreased. Smaller spheres with a snowflake-like shape were formed. Finally, large micrometric "spherical-type objects" with a hole were obtained by aggregation of these small snowflakes. In fact, micrometric spheres of MnO 2 have been frequently reported to form upon solvothermal synthesis [59].

For the highest Mn:Ce ratio of 1, the Mn-Ce solid solution existed together with the hausmannite phase (Mn 3 O 4 ). The snowflake like morphology was maintained but the specific micrometric sphere morphology was lost. In the case of a pure manganese sample, flower-like micrometric Mn 3 O 4 objects were obtained.

The evolution of textural properties summarized in the Table 1 might be better understood in the light of the TEM observations. In the case of pure ceria, the aggregation of large rectangular crystallites induced the formation of large pores. Indeed, the "orientated aggregation" of small ceria crystallites into larger irregular truncated octahedral crystals was shown to induce the formation of pores [9]. Moreover, in the CM 0.25 P 3 sample, smaller and more isotropic crystallites were aggregated, in a more compact stacking, leading to the absence of any porosity. Finally, in the case of the CM 0.5 P 3 sample, the snowflake-like porous structure was obtained by the aggregation of very small crystallites, which explained the higher specific surface area and the higher porous volume.

EPR results showed that there was some Mn II inside the ceria lattice in substitution position.

The Mn 3 O 4 phase was detected by XRD but it could not be evidenced by EPR. Indeed, strongly interacting Mn III species are difficult to observe and strongly interacting Mn II might not be distinguished from the isolated Mn II species [55]. Manganese IV was also evidenced, either as Mn IV inserted in the ceria lattice or as Mn IV as the outer ceria surface. These data should be compared with the literature. At least two types of Mn II species have already been observed in some Mn-Ce oxide samples. The absence of Mn IV species was reported. The presence of Mn III species, difficult to prove by EPR analysis, was deduced from other experiments [37]. In other studies, it has been clearly observed that Mn II replaced Ce IV in ceria, as evidenced by XPS [60] or Rietveld analysis [61].

Two recent studies have shown that a limited fraction of manganese was present in a solid solution with ceria, while the other part of manganese was segregated around the ceria crystallites in the grain boundary. It was manifested by a crystallite size decrease and a specific surface area increase [43,62]. Our EPR results were also consistent with such hypothesis since the Mn IV signal could be explained by the segregation of manganese at the outer ceria surface. In fact, such kind of segregation at the grain boundary is frequently observed when ceria is doped with rare-earths [63,64] and the accumulation of cation at the grain boundaries is found to coincide with the oxygen vacancy accumulation. The micron-sized spherical morphology obtained in this work for the Mn:Ce ratio of 0.5, may result from this manganese segregation, inducing a cohesion between the Mn-Ce oxide crystallites. Our work could be compared with previous studies on mixed manganese cerium oxide via solvothermal synthesis. A new snowflake-like porous morphology in the CeMn oxide system, formed by the oriented attachment of small crystallites, has been obtained by the proposed method, which differs from the rods and cubes or flowers previously synthesized [24,25]. The particular snow-flake porous morphology of the cerium-manganese mixed oxide which has been obtained in this work results from the polyvinylpyrrolidone addition. Indeed, the PVP polymer behaves as a capping agent around the ceria seeds and promotes the manganese ions dispersion on the ceria seeds.

Conclusions

Reactions involving manganese and cerium salts in the presence of PVP as a capping ligand and structuring agent in a solvothermal process led to Mn-Ce oxides with snowflake-like morphology.

Manganese was partly introduced as Mn II inside the ceria lattice. For a Mn:Ce ratio of 0.25, small spheres of 100 nm in diameter were formed, consisting of a Mn-Ce oxide solid solution. The other fraction of the manganese was present as Mn IV, near the ceria surface. Furthermore, these later crystallites were found to be highly oriented. For a Mn:Ce ratio of 0.5, smaller snowflake-like spheres of a MnCe oxide were successfully synthesized. These snowflakes were aggregated into large porous micrometric spheres. At higher Mn:Ce ratio, snowflake-like spheres of Mn-Ce oxide were mixed with Mn 3 O 4 . The hydrothermal stability of such compound was also studied and the crystallization of a bixbyite phase (Mn 2 O 3 ), out of the ceria was observed. However, the texture of the solid was preserved.

  2 revealed (i) a dehydration step characterized by a first weight loss related to the physically adsorbed water; (ii) a second weight loss, attributed to the reduction of MnO 2 to Mn 5 O 8 (Fig. 9a); (iii) a third weight loss, starting at around 400°C, associated to an endothermic peak on SDTA curve, and corresponding to the formation of Mn 2 O 3 ; (iv) a fourth weight loss, at around 900°C, related to the subsequent endothermic reduction of Mn 2 O 3 to α-Mn 3 O 4 and (v) on the SDTA curve, a last endothermic transition at about 1140°C due to the transformation of α-Mn 3 O 4 into β-Mn 3 O 4 . TGA of CM 0 P 3 (Fig. 9b) showed a first weight loss below 200°C attributed to the loss of water and a second continuous weight loss from

  Figures caption.

Figure 1 :

 1 Figure 1: a, b: TEM pictures of the CM 0.5 P 3 sample as obtained after the solvothermal treatment at 200°C.

Figure 2 :

 2 Figure 2: a: TGA and b: SDTA curves of the CM 0.5 P 3 sample after the solvothermal treatment at 200°C.

Figure 3 :

 3 Figure 3: X-ray diffraction patterns of the CM x P y samples. a: CM 0 P 3 , CM 0.25 P 3 , CM 0.5 P 3 , CM 1 P 3 , MP 3 samples, Inset: zoom on the (111) reflection; b: CM 0.5 P 0 , CM 0.5 P 1 , CM 0.5 P 3 samples, Inset: zoom on the (111) reflection.

Figure 4 :

 4 Figure 4: SEM pictures of the CM x P y samples. a: CM 0.5 P 0 ; b: CM 0.5 P 1 ; c: CM 0.5 P 3 ; d: CM 0 P 3 ; e: CM 0.25 P 3 ; f: MP 3 .

Figure 5 :

 5 Figure 5: TEM pictures of the CM x P y samples. a: CM 0.5 P 3 ; b: CM 0.25 P 3 ; c: CM 0.5 P 3 ; d: CM 1 P 3; e: CM 0.5 P 0; f: CM 0.5 P 1 .

Figure 6 :

 6 Figure 6: TEM, HRTEM pictures and electron diffraction patterns on a single sphere. a, b c: CM 0 P 3 ; d, e, f: CM 0.25 P 3 ; g, h, i: CM 0.5 P 3 samples.

Figure 7 a

 7 Figure 7 a : EPR spectra, parallel mode, CM 0.5 P 3 (15K), CM 0.5 P 0 (15K), CM 0.25 P 3 (15K), CM 0.5 P 3 (25K); b : EPR spectra, perpendicular mode, CM 0.5 P 3 (15K), CM 0.5 P 0 (15K), CM 0.25 P 3 (15K), CM 0.5 P 3 (25K); c : EPR spectra, parallel mode, CM 0.5 P 3 (25K), experimental, simulated.

Figure 8 :

 8 Figure 8: Raman spectra of the CM x P y samples. a: CM 0 P 0 ; b: CM 0 P 3 ; c: CM 0.5 P 0 ; d: CM 0.5 P 3 . e: CM 1 P 0 ; f: CM 1 P 3 .

Figure 9 :

 9 Figure 9: TGA/SDTA curves of the CM x P 3 samples. a: MnO 2 reference sample; b: CM 0 P 3 sample; c: CM 0.25 P 0 , CM 0.5 P 3 and CM 1 P 3 samples.

Figure 10 :

 10 Figure 10: Nitrogen adsorption/desorption isotherms and pore size distribution from BJH in insert of the CM x P 3 samples: a: CM 0.5 P 3 ; b: CM 0.5 P 3 -h; c: CM 0.5 P 0 ; d: CM 0.5 P 1 .

Figure 11 :

 11 Figure 11: a: X-ray diffraction patterns of the samples before: CM 0.5 P 3 and after: CM 0.5 P 3-h, hydrothermal treatment; Zoom on the (111) ceria reflexion in insert. b: SEM picture of the CM 0.5 P 3-h sample. c, d: TEM pictures of the CM 0.5 P 3-h sample.

Table 1 :

 1 Textural properties of CM x P y samples.

	Table:
		Mn/Ce	Specific	Mesoporous	Porous	Max pore	(220)	Crystallite
		ratio	Surface	volume	volume	radius	width (°)	size (nm)
			area (m²/g)	(cm	3 /g)	(cm	3 /g)	(nm)
	CM 0 P 3	0/1	59	0.10	0.22	7	0.55	16
	CM 0.25 P 3	0.25/1	58	0.03	0.06	1.7	1.60	5.8
	CM 0.5 P 3	0.5/1	108	0.16	0.21	11	2.33	3.9
	CM 1 P 3	1/1	70	0.05	0.06	2.0	2.08	4.0
	CM 0.5 P 0	0.5/1	30	0.03	0.05	1.7	1.36	6.8
	CM 0.5 P 1	0.5/1	97	0.16	0.19	1.9	1.68	5.5
	CM 0.5 P 3 -h	0.5/1	92	0.21	0.25	14	2.17	4.2

Acknowlegments

This work was financially supported by the Région Rhône-Alpes. The authors gratefully acknowledge the Region Rhône-Alpes for the Ph. D. grant awarded to France Schmit for the period 2011-2014. The authors gratefully acknowledge the national EPR network TGE RENARD, FR 3443-CNRS and the CTµ platform of electronic microscopy (University Lyon 1). Nicolas Blanchard, Institut Lumière Matière, Université Lyon 1, is acknowledged for his help during the TEM observations.


Jingyuan Linda Zhang

Hitoshi Ishiwata

Thomas M Babinec

Marina Radulaski

Kai Müller

Konstantinos G Lagoudakis

Constantin Dory

Jeremy Dahl

Veronique Soulière

Gabriel Ferro

Andrey A Fokin

Peter R Schreiner

Zhi-Xun Shen

Nicholas A Melosh

Jelena Vučković

Hybrid Group IV Nanophotonic Structures Incorporating Diamond Silicon-Vacancy Color Centers

Keywords: diamond, nanodiamonds, nanophotonics, nanofabrication, silicon carbide, siliconvacancy (SiV) color center in diamond

We demonstrate a new approach for engineering group IV semiconductor-based quantum photonic structures containing negatively charged silicon-vacancy (SiV -) color centers in diamond as quantum emitters. Hybrid SiC/diamond structures are realized by combining the growth of nanoand micro-diamonds on silicon carbide (3C or 4H polytype) substrates, with the subsequent use of these diamond crystals as a hard mask for pattern transfer. SiV -color centers are incorporated in diamond during its synthesis from molecular diamond seeds (diamondoids), with no need for ionimplantation or annealing. We show that the same growth technique can be used to grow a diamond layer controllably doped with SiV -on top of a high purity bulk diamond, in which we subsequently fabricate nanopillar arrays containing high quality SiV -centers. Scanning confocal photoluminescence measurements reveal optically active SiV -lines both at room temperature and low temperature (5 K) from all fabricated structures, and, in particular, very narrow linewidths and small inhomogeneous broadening of SiV -lines from all-diamond nano-pillar arrays, which is a critical requirement for quantum computation. At low temperatures (5 K) we observe in these structures the signature typical of SiV -centers in bulk diamond, consistent with a double lambda.

These results indicate that high quality color centers can be incorporated into nanophotonic structures synthetically with properties equivalent to those in bulk diamond, thereby opening opportunities for applications in classical and quantum information processing.

TABLE OF CONTENT GRAPHIC TEXT

Quantum emitters in solids are an important resource for diverse applications including quantum information processing (QIP) and sensing. [1][2][3] Some of the most common approaches include emitters such as InAs quantum dots and nitrogen-vacancy (NV -) color centers that are incorporated into devices composed of a single material such as GaAs and diamond, respectively. 4,5 However, these approaches suffer from drawbacks including the limited coherence times of spins in InAs/GaAs quantum dots, 6 the necessity for operation at cryogenic temperatures, 4 incompatibility with Si CMOS platforms (as is needed for devices in optical interconnects), and large inhomogeneous broadening. 7 On the other hand, for NV -centers in diamond, drawbacks include a lack of a second order optical nonlinearity needed for efficient frequency conversion, 8 degradation of NV -center properties resulting from proximity of etched surfaces, and nonstandard fabrication methods needed to carve optical structures in bulk diamond. 9,10 Hybrid solid- state schemes consisting of quantum emitters in individual nanocrystals coupled to semiconductor optical cavities in other materials have also been demonstrated; 11 however, they typically require AFM 'pick-and-place' assembly techniques 12 or complex film transfers. [13][14][15][16] Silicon-vacancy (SiV -) color centers in diamond have recently emerged as promising candidates for QIP applications. 17 Due to the inversion symmetry of the SiV -centers, they display superior spectral stability and narrow inhomogeneous broadening, with room temperature linewidths of a few nanometers and nearly transform-limited linewidths at low temperature. 17 Even more importantly, 70% of the photons are emitted into the strong zero-phonon line, compared to only 3% for NV -centers. Single photon generation from SiV -centers has by now been routinely demonstrated, 18,19 and their application in QIP, such as quantum key distribution (QKD), has been considered. 20 Moreover, this emitter offers a Λ-system energy level configuration, which is needed for all optical electron spin control, in weak or zero magnetic fields. This contrasts with the InAs/GaAs QDs that have to be charged with an extra electron or hole and placed in a strong magnetic field to realize a Λ-system. 21 Recently, optical initialization, readout, and coherent preparation of the electron spin of the SiV -center in bulk diamond was demonstrated. 22,23 However, as opposed to InAs/GaAs quantum dots, SiV -s have yet to be incorporated into high quality nanophotonic structures without affecting their spectral signatures, 24 which will enable strong interaction and efficient interfacing with photons. The orbital relaxation limited spin coherence time of SiV -centers has been measured to be in the 10s of nanosecond range, 22,23 which is not yet competitive with the spin coherence time in charged InGaAs quantum dots. 25 However, there may be a possibility to improve the coherence by reducing the phonon density of states in the SiV - environment 26 via engineering nanostructures similar to those shown in this article.

Here we report a new approach for implementing nanophotonic structures with embedded diamond SiV -centers of high quality, and apply it to two types of devices: diamond-SiC hybrid structures and all-diamond nanopillars arrays. Nanodiamonds and diamond films doped with SiV - centers are grown from molecular diamond seeds (diamondoids) on SiC and high purity bulk diamond substrates, respectively. SiV -color centers are incorporated in diamond during its synthesis, with no need for ion-implantation or annealing. Nanodiamonds grown from diamondoids in our approach are devoid of nitrogen vacancy centers, which could be present in the Ultra Dispersed Diamonds (UDD). 27 For hybrid diamond-SiC structures, functionalized nanodiamonds are used as a hard mask for pattern transfer into the 3C and 4H-SiC substrates, followed by subsequent removal of a buried Si sacrificial layer in the case of 3C-SiC. This simple technique results in high yield of devices containing SiV -centers, avoiding the AFM 'pick-andplace' or film transfers. Moreover, combining two group IV semiconductors -diamond and silicon carbide -could potentially address challenges in each individual material, including difficulty in fabricating planar photonic structures in diamond, lack of doping for electrical devices, absence of second order optical nonlinearity in diamond, [START_REF] Zaitzev | Optical Properties of Diamond[END_REF][START_REF] Harris | [END_REF][30][31] and less developed quantum emitters in SiC. 32,33 In addition, diamond-SiC devices could benefit from the mature fabrication methods developed for SiC, which has produced high quality factor optical microcavities. 30,31 Finally, the method we describe in the article allows for positioning of diamond relative to SiC structures, as diamond itself is used as an etch mask. On the other hand, all-diamond monolithic nanopillar arrays are nanofabricated in SiV -doped diamond films grown on bulk diamond substrates, using standard microfabrication methods (lithography and dry etching). The resulting nanopillars contain high quality SiV -centers with very narrow linewidths and very low inhomogeneous broadening, thereby enabling implementation of quantum photonic devices containing identical quantum emitters, as needed for quantum simulations 34 and quantum networks. 35 As explained below, our growth methods allow us to control vertical positioning of SiV -inside diamond films to ~10nm thickness at a chosen depth.

Growth and fabrication of structures

High quality nanodiamond crystals and diamond films are CVD grown on SiC and bulk diamond substrates, respectively, starting from diamondoid seeds covalently attached to a SiC or diamond surface (Fig. 1(a)). Diamondoids are a class of face-fused adamantane (C 10 H 16 ) building blocks whose extension eventually leads to macroscopic diamond. [36][37][38] The lower diamondoids are adamantane (C 10 H 16 ), diamantane (C 14 H 20 ) and triamantane (C 18 H 24 ), while the higher diamondoids begin with isomeric tetramantane (C 22 H 28 ) and pentamantane (C 26 H 32 ). While the lower diamondoids can readily be synthesized, 37,38 the higher diamondoids essentially are only accessible from petroleum sources via high pressure liquid chromatography (HPLC). 39 Unlike typical detonation diamond used to seed diamond CVD growth, diamondoids are free from nitrogen and graphitic impurities, with a precisely known molecular structure like most other small organic molecules. Diamondoids have been applied to diamond growth, [40][41][42] electron imaging 43 and electron emission devices. 44 Here, we covalently bond 7-dichlorophosphoryl[1(2,3)4]pentamantane 45 as a seed for high quality growth of fluorescent diamond nanoparticles and to form a bond between heteroepitaxial diamond layer and substrate.

The growth of the diamond is illustrated in Figure 1 (b). First, an oxide layer is generated on the substrate with exposure to oxygen plasma for 5 min at 400 mTorr pressure and 100 W power.

For the devices presented in this paper, bulk 4H-SiC wafers (Cree) as well as heteroepitaxial 3C-SiC(100) thin films (~150 nm) grown on Si (100) via a standard two-step procedure were used. 46 The in the as-grown nanodiamonds. They were incorporated during the 'growth step' through diffusion of Si atoms from the plasma etching of SiC substrate, without need for subsequent annealing or ion implantation. Our growth approach does not degrade the quality of the transition lines of the SiV -center, and leads to properties comparable to SiV -in bulk diamond. 17 Although we don't have control over lateral positioning of SiV -centers in our growth approach, as opposed to ion implantation, we do have control over their vertical positioning, which enables us to grow SiV -s only within ~10 nm thick diamond layer at a chosen depth.

The process flow for fabricating hybrid diamond-SiC devices is shown in Figure 2. First, high quality nano-and micro-diamond doped with SiV -was grown on 3C-and 4H-SiC as described above. The high quality of the seeding and growth process ensures optical quality of the fabricated structures, as will be demonstrated in the PL measurements below. Then, the diamond particle shape was transferred into the substrate through anisotropic etching of the substrate using the diamond crystals as the etch mask. Lastly, any additional processing (e.g., undercutting in case of 3C-SiC on Si substrate, or thinning in case of 4H-SiC) can be performed. As a proof of concept, we have applied this approach to generate two different types of hybrid diamond-SiC nanophotonic structures: (i) diamond -4H-SiC nanowires (also applicable to 3C-SiC) and (ii) diamond-3C-SiC hemispherical microdome (akin to whispering gallery mode resonators). For the diamond -4H-SiC nanowires, approximately 500 nm diameter diamond nanocrystals were first grown on 4H-SiC, and then used as hard mask to etch the underlying substrate using inductively coupled plasma (ICP) etching with HBr/Cl 2 chemistry, as illustrated in Figure 2(a). The SEM image of a diamond -4H-SiC nanowires is shown in Figure 2(b). Alternatively, for fabricating the diamond-3C-SiC microdome structures, diamond microcrystals with diameters ~ 2 μm were first grown on 150 nm thick 3C-SiC epitaxial film on Si. 46 Then, the pattern of randomly distributed microdiamonds was transferred through 3C-SiC using the same etch recipe as above, followed by undercutting the sacrificial Si using the XeF 2 vapor phase silicon etcher, as illustrated in Figure 2 For monolithic nanopillar arrays, we grow a homoepitaxial layer of SiV -center doped thin film grown on diamond substrate using the same method as described earlier for the nanodiamonds. Starting with a bulk diamond substrate (Element Six, type Ib), a 70 nm diamond film containing SiV -was grown homoepitaxially via microwave plasma-assisted chemical vapor deposition (MPCVD). The SiV -centers are introduced through diffusion of Si atoms from plasma etching of the SiC placed near the diamond substrate during the growth. After the film growth, the nanopillars arrays are defined lithographically using evaporated gold as hard mask, as illustrated in Supplementary Figure 1(a). Figures 2(f) and (g) show SEM images of a typical fabricated array and an individual nanopillar.

Optical characterization

The presence of SiV -centers in all described structures is confirmed by scanning confocal microscopy measurements. The custom made laser scanning confocal microscope consists of 532 nm continuous wave (CW) pump laser focused onto the sample through a high numerical aperture NA = 0.75 microscope objective, as shown in Figure 3(a). The photoluminescence (PL) is collected through the same objective and is sent into a single mode collection fiber with a dichroic mirror.

A scanning galvanometer in the common path of the pump and collection scans the focal spot across the sample surface. This allows producing PL maps of the sample, as well as addressing individual devices using the scanning mirror. The collected emission is directed onto an avalanche photodiode (APD) for generating the PL map, or a high-resolution spectrometer for spectral characterization.

A typical laser scanning confocal microscope image for hybrid diamond-SiC nanowires is shown in Figure 3(b) (similar PL maps were obtained for nanodiamonds as well as SiC microdomes). The bright areas with high count rates indicate the presence of randomly distributed nanodiamonds containing SiV -centers in described photonic structures. The room temperature PL spectrum of a typical diamond-SiC nanowire is presented in Figure 3(c) and exhibits a narrow emission peak at 738 nm, corresponding to the emission from SiV -centers embedded in the diamond. 17 Similar spectra were observed prior to the fabrication of the photonic structures.

Therefore, this confirms that the SiV -centers in the nano-and micro-diamond crystals before the fabrication process are retained in the photonic structures post-fabrication. The low temperature PL spectra of SiV -s in nanodiamonds grown on SiC can be seen in Supplementary Figure 2. In this hybrid system at low temperature, SiV -s feature multiple emitter lines and strain induced spectral shifts in transitions, likely resulting from the lattice mismatch between the SiC substrate and diamond grown on top. Growth of such hybrid systems on substrates that have a better lattice match to diamond could be a route to addressing this issue, or simply a growth of all diamond homoepitaxial systems presented below.

In addition, a laser scanning confocal microscopy map of the diamond nanopillar array sample is shown in Figure 4 Therefore, in addition to narrow linewidths, the demonstrated SiV -centers incorporated in diamond nanopillar structures also feature small inhomogeneous broadening, potentially arising from the lower defect density and lower strain in the homoepitaxial diamond nanopillars (see Supporting Information). Although we do not observe inhomogeneous broadening of SiV -lines at low temperature inside 250 nm alldiamond homoepitaxial nanopillars (within our spectrometer resolution limit of 4 GHz), we start to see some splitting in smaller, 130 nm diameter nanopillars (see Supplementary Figure 3). By counting the number of line quadruplets inside such structures we can estimate that on average we have ~5 SiV -s inside a 130 nm nanopillar, or doping density of 5×10 15 cm -3 . This agrees with our estimate of the density of SiV -s in hybrid diamond-SiC structures doped at similar level. Therefore, reduction in doping density is needed to isolate a single SiV -inside such structures. As an additional confirmation of the fact that we have multiple SiV -s in our current structures, we have performed g (2) (t) measurements on the same 130 nm pillars shown in the Supplementary Figure 3.

Antibunching is absent, as expected. The lifetime measurements on nanopillars reveals a mean lifetime of 1.18 ns (see Supplementary Figure 4), and both the PL linewidth and lifetime are comparable to that of single SiV -centers in bulk. 17 The small inhomogeneous broadening in the larger pillars suggests significant potential for entangling different quantum emitters on a chip via photon interference, and for using such arrays to generate large entangled photon states efficiently -as needed for quantum networks and quantum simulations. 34,35 Therefore, our approach for engineering quantum photonic structures incorporating diamond SiV -not only enables the implementation of an efficient interface between a quantum emitter and a photon, but also does so without degradation of the properties of the quantum emitter itself. Moreover, the demonstrated nanopillar structures may also offer the opportunity to prolong the spin coherence time via reduced phonon density of states, thereby improving the potential of this system as a solid state quantum memory.

In conclusion, we have demonstrated a new approach for implementing nanophotonic structures with selectively incorporated high quality SiV -centers in nanodiamonds or diamond films during MPCVD growth (without need for ion implementation or annealing). We have successfully applied this approach to two types of structures: hybrid diamond-SiC structures (SiC microdomes and nanowires with diamond tips) and all-diamond nanopillar arrays with quantum emitters at tips of pillars. The hybrid diamond-SiC devices are fabricated by pattern transfer into the SiC substrate using diamond nanocrystal as hard mask, which simplifies fabrication, but also retains the SiV -centers post fabrication, as illustrated by the room temperature spectra. On the other hand, monolithic diamond nanopillar arrays were fabricated through a diamond thin film grown on high purity diamond substrate, followed by conventional microfabrication. In particular, we note the narrow linewidths, absence of bleaching, and small inhomogeneous broadening in the diamond nanopillars array at low temperature. This can open opportunities for entangling multiple quantum emitters on the chip. Future characterization will involve resonance fluorescence spectroscopy and Hong-Ou-Mandel interferometry. The developed structures could be employed as large arrays of individually addressable quantum emitters of indistinguishable photons for use in quantum networks. Future directions could also encompass utilizing the reduced phonon density of states to prolong the spin coherence time, in combination with the low strain environment, to improve the potential of this system as a solid state quantum memory. To fabricate the nanopillars arrays, 5 nm HfO 2 was deposited as an adhesion layer via Atomic Layer Deposition (ALD) following the diamond film growth. Arrays of nanopillars with diameter 115-250 nm diameter and 150 nm height were defined using electron-beam lithography followed by ICP RIE, with electron-beam evaporated gold as hard mask. Lastly, the gold hard mask and adhesion layer were removed using gold etch and piranha clean recipes, leaving nanopillars with silicon vacancy centers near the top of the nanopillars.

Low temperature PL of hybrid diamond-SiC systems

Supplementary Figure 2: (Left) Low temperature PL emission spectrum of a 200 nm nanodiamond on 4H-SiC. (Right) Low temperature PL emission spectrum of a diamond-SiC nanowire as described in the manuscript. In this hybrid system at low temperature, SiV -s feature multiple emitter lines and strain induced spectral shifts in transitions, likely resulting from the lattice mismatch between the SiC substrate and diamond grown on top. As can be seen from the low temperature spectra of nanodiamonds on SiC (Figure S2), SiV -s in these hybrid systems feature multiple emitter lines and strain induced spectral shifting at low temperature. The left figure shows the low temperature PL spectrum of a 200 nm nanodiamond on 4H-SiC, while the right figure shows the low temperature PL spectrum of the diamond-SiC nanowires in the manuscript. Therefore, homoepitaxial system results in much smaller strain and smaller inhomogenous broadening.

Doping density

Supplementary Figure 3: Figure R2: (Left) Low temperature PL emission spectrum of a nanopillar 130 nm in diameter. (Right) Representative g (2) (t) measurements on the smallest nanopillars produced with ~130 nm diameter. g (2) (0)~0.81 is estimated from the fit, indicating multiple emitters inside the pillar, as expected from the number of line quadruplets shown in the left figure. Since we used the same Si doping densities both for the heteroepitaxial and homoepitaxial growth process, we can use strain split spectra of heteroepitaxial structures to estimate the number of SiV -s per pillar. The shown spectra of diamond grown on SiC suggest multiple (2+) emitters per nanodiamond of diameter 200 nm. In the nanopillar structures, we see no strain splitting in 250 nm diameter pillars, as shown in the main text, but strain splitting starts to manifest in 130 nm pillars, as shown in Supplementary Figure 3(left). One emitter per pillar (~ 130 nm diameter) corresponds to a SiV -density of 1.0×10 15 cm -3 , so we expect an emitter density of several 10 15 cm - 3 . This is consistent with the density estimated from the ~200 nm diameter nanodiamond.

We have then performed g (2) (t) measurements on the all diamond, homoepitaxial nanopillars with the smallest diameters (~ 130nm diameter), as shown in Supplementary Figure 3(right), but didn't observe anti-bunching. This is in agreement with estimated SiV -density above.

However, multiple emitters with small inhomogeneous broadening embedded in photonic structures are also very interesting for multi-emitter CQED, which is not possible with quantum dots (as a result of large inhomogeneous broadening). This is in fact another topic of our research, as in such multi-emitter system one can increase Rabi splitting in proportion to √N, where N is the number of emitters. This increases the operating speed of the devices based on the strong coupling CQED effects, and makes reaching the strong coupling regime easier.

  (c). The SEM images of diamond -3C-SiC microdomes are shown in Figure 2(d) and (e).

  (a), revealing strong PL at the locations of the nanopillars. High resolution spectra of the nanopillars (Figure 4(b)) confirm PL from SiV -at room temperature. At low temperature (~5 K) the spectra (Figure 4(c) and (d)) show four distinct lines, consistent with the signature of the electronic energy levels of unstrained SiV -centers in diamond -and so far observed in bulk diamond only. 17 Repeatable measurements of the spectrum of nanopillar A (as highlighted in Figure 4(a)), are presented in Figure 4(c) and clearly demonstrate the absence of bleaching over minutes timescale. Furthermore, different nanopillars in the same array (i.e., pillars B and C) show strong spectral overlap, as shown in Figure 4(d).

Figure 1 .

 1 Figure 1. (a) Molecular structure of 7-dichlorophosphoryl[1(2,3)4]pentamantane. (b) Schematic of 7-dichlorophosphoryl[1(2,3)4]pentamantane on an oxide layer formed on top of SiC substrate. (c) Scanning electron micrograph (SEM) of 500 nm diameter nanodiamonds grown on 4H-SiC substrate. (d) SEM of a micrometer size diamond on 3C-SiC substrate.

Figure 2 .Figure 3 .Figure 4 .

 234 Figure 2. (a) Process flow for fabricating diamond -SiC nanowires through hard mask pattern transfer. The red dots represent the silicon vacancy centers in the diamond nanocrystals. (b) SEM image of a typical nanowire. (c) Process flow for hybrid diamond -SiC microdome structures fabricated through hard mask pattern transfer. (d) SEM image of a typical microdome structure. The high-quality diamond, as well as the heteroepitaxial interface are visible. (e) SEM image of an ensemble of microdome structures. (f) SEM image of a fabricated monolithic nanopillar array and (g) close-up SEM image of a monolithic nanopillar.

  ]pentamantane. This process results in the generation of a covalently attached [1(2,3)4]pentamantane monolayer on the silicon carbide samples. From here, the sample is placed in a CVD reactor for the 'nucleation step' (gas mixture H 2 : 5 sccm, CH 4 : 10 sccm, Ar: 90 sccm, substrate temperature: 450 °C, microwave power: 300 W, pressure: 23 Torr) for ~20 min to enhance nucleation density observed from diamondoids. After nucleation, a 'growth step' (gas mixture H 2 : 300 sccm CH 4 : 3-7.5 sccm, substrate temperature: 830 °C, microwave power: 1300

	sample	is	then	soaked	in	toluene	solution	containing	1mM	7-
	dichlorophosphoryl[1(2,3)4

W, pressure: 30 Torr, 1-2.5% CH 4 in H 2 carrier gas) is performed, with growth times calibrated to desired nanoparticle size. High quality diamond crystals with grain sizes ranging from 500 nm to 2 µm can be seen in the scanning electron microscope (SEM) images in Figures

1(c) and (d

). As will be discussed below, in photoluminescence (PL) experiments we observe the presence of SiV -

ACKNOWLEDGMENT Financial support for materials synthesis is provided by the DOE Office of Basic Energy Sciences, Division of Materials Sciences through Stanford Institute for Materials and Energy Sciences (SIMES) under contract DE-AC02-76SF00515. Financial support for this work is also provided by Stanford Institute for Materials and Energy Sciences (SIMES), National Science Foundation grant DMR-1406028, AFOSR MURI on quantum metaphotonics and metamaterials, and the Deutsche Forschungsgemeinschaft (Schr 597/23-1). This work was performed in part at the Stanford Nanofabrication Facility of NNIN supported by the National Science Foundation under Grant No. ECS-9731293, and Stanford Nano Shared Facility. JLZ acknowledges support from the Stanford Graduate Fellowship. KM acknowledges support from the Alexander von Humboldt Foundation. LJZ thanks Yan-Kai Tzeng for the helpful discussions.

ASSOCIATED CONTENT Supporting Information.

Process flow for fabricating diamond nanopillar arrays, low temperature PL of hybrid diamond-SiC systems, and spectral, lifetime, and g (2) (t) measurements on all-diamond nanopillars. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION

Corresponding Author

* Email: Jelena Vučković: jela@stanford.edu, Nicholas A. Melosh: nmelosh@stanford.edu, and Zhi-Xun Shen: zxshen@stanford.edu.

Present Addresses

Thomas M. Babinec's present address: US Army Research Laboratory, 2800 Powder Mill Road, Adelphi MD 20783

Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ‡JLZ and HI contributed equally.

Notes

The authors declare no competing financial interest.

Supporting Information

Fabrication of nanopillar arrays

Supplementary Figure 1. (a) Process flow for fabricating diamond nanopillar arrays. Diamond film containing SiV -is grown homoepitaxially on bulk diamond substrate (Element Six, type Ib). Subsequent ALD deposition of 5 nm HfO 2 facilitates metal adhesion. The nanopillar array is defined by electron-beam lithography in bilayer positive resist PMMA, followed by evaporation of gold as hard mask and the lift-off of PMMA. Lastly, the nanopillar array pattern is transferred into the diamond substrate past the SiV -layer so that the background is SiV -free, and the gold and HfO The linewidth of all presented PL measurements was limited by the spectrometer resolution of 4.04 GHz, and is therefore greater than or equal to the true inhomogeneously broadened linewidth measured with other methods such as photoluminescence excitation.

We measured the spectra and lifetimes at the same spots in nanopillars (diameter ~ 130 -150 nm), an example of which is shown in Supplementary Figure 4 (left, center). In the left figure, the spectrum is taken on a nanopillar with 150 nm diameter, while the strongest transition has a FWHM of 28.3 GHz (or 0.0513 nm, 117 µeV). At the same spot, the lifetime was measured to be noted that we see no strain induced spectral shifting in 250 nm pillars, but we did start to see some spectral shifting in 130 nm pillars, as shown in Supplementary Figure 4(right).