A Takacs

H Aubert

A Luca

S Charlot

S Fredon

L Despoisse

Rectenna Design for K Band Application

Keywords: Rectenna, energy harvesting, satellite application I

An ultra-compact (2.5 cm 2 ) rectenna design based on an innovative topology using a coplanar stripline connected to a crossed dipoles antenna array is reported here. Experimental results in the K-band show that a DC power up to 0.9 mW can be harvested on a 0.3 kΩ load by using the proposed compact rectenna for incident power density of 1413 µW/cm 2 . The targeted application for this rectenna is to power autonomous wireless sensors for satellite health monitoring.

INTRODUCTION

Recently, the topics of energy harvesting for satellite health monitoring was explored [START_REF] Takacs | Microwave energy harvesting for satellite applications[END_REF]- [START_REF] Takacs | Compact Rectenna for Space Application[END_REF]. The idea consists of harvesting the spill-over loss of microwave antenna (operating in C, X, Ku, K or Ka bands) of broadcasting satellites by using appropriate rectenna designs. This harvested DC power available over the satellite lifetime can be used to power autonomous wireless sensors for satellite health monitoring. Rectenna topologies working on frequencies beyond 10 GHz were addressed for microwave power transmission [START_REF] Strassner | Microwave Power Transmission: Historical Milestones and system Components[END_REF]- [START_REF] Ren | 35 GHz rectifying antenna for wireless power transmission[END_REF], 24 GHz ISM band applications [START_REF] Shinohara | Development of 24 GHz rectennas for Fixed Wireless Access[END_REF]- [START_REF] Collado | 24 GHz Substrate Integrated Waveguide (SIW) Rectenna for Energy Harvesting and Wireless Power Transmission[END_REF] or more recently satellite hearth monitoring [START_REF] Takacs | Microwave energy harvesting for satellite applications[END_REF]- [START_REF] Takacs | Compact Rectenna for Space Application[END_REF]. This paper addresses an innovative rectenna topology for K band application. The design methodology, simulation and experimental results are presented in order to underline the performances of the proposed design.

II. RECTENNA DESIGN

A. Topology

The proposed topology is based on the use of a high gain planar antenna directly connected with a high frequency rectifier using (only) one Schottky diode. Fig. 1 shows the top view of the rectenna while its bottom view is represented in Fig. 2. The antenna consists of an array of two planar crossed dipoles. The rectifier consists of (i) a Schottky diode mounted on the top side of the planar circuit board (PCB) and (ii) a lumped filtering capacitor of 1.5 pF and a load (see Fig. 2). Thus, the low-pass filtering function is achieved by the coplanar strip supported structure mounted on the bottom side of the rectenna. The connection between the diode and the lowpass filter is performed by using two via holes. The matching between antenna and rectifier input impedances is achieved by properly controlling the mounting position of the Schottky diode. Rectennas using coplanar stripline structures are reported in [START_REF] Takacs | Microwave Power Harvesting for Satellite Health Monitoring[END_REF], [START_REF] Takacs | Compact Rectenna for Space Application[END_REF] and [START_REF] Yoo | Theoretical and experimental development of 10 and 35 GHz rectennas[END_REF] while microstrip supported topologies are reported in [START_REF] Takacs | Microwave Power Harvesting for Satellite Health Monitoring[END_REF]- [START_REF] Takacs | Compact Rectenna for Space Application[END_REF], [START_REF] Ren | 35 GHz rectifying antenna for wireless power transmission[END_REF]- [START_REF] Ladan | High Efficiency Low-Power Microwave Rectifier for Wireless Energy Harvesting[END_REF]. A substrate integrated waveguide (SIW) rectenna is reported in [START_REF] Collado | 24 GHz Substrate Integrated Waveguide (SIW) Rectenna for Energy Harvesting and Wireless Power Transmission[END_REF].

B. Antenna design

The antenna located on the top side can be viewed as an array composed by 4 printed half-wavelength dipoles. This antenna is called here the crossed dipoles array antenna (CDAA). The maximum theoretical gain (Gmax in dBi) of such an array antenna can be computed using the following formula:

Gmax=10•log(4)+Gd=6+2.15=8.15 dBi (1)
where Gd=2.15 dBi is the gain of a lossless half-wavelength dipole. This gain is 3 dB higher when a metallic reflector is properly positioned below the antenna surface. Thus the expected maximum gain is 11.15 dBi. The main geometrical parameters for optimizing CDAA performances are: the length of the crossed dipole arms (ld), the angle between crossed dipoles (α), the array step (d), the strip width (w) and the gap between the strips of the coplanar stripline (g). The length of the crossed dipole arms (ld) and the array step must be close to the half-wavelength. The optimization procedure for CDAA is described in the next section. A standard 3D Cartesian coordinate system with the vertical Oz axis perpendicular to the rectenna/CDAA surface is chosen here. The diode was replaced by a voltage port while the capacitance (1.5 pF) and the load (e.g. 300 Ω) were modeled as port loads. By adopting this simulation model the impact on antenna performances of the overall rectifier layout (except the non-linear behavior of the diode) is estimated. Fig. 3 shows the simulated gain (by using FEKO software) in two vertical planes (xOz-plane: phi=0° and yOz-plane: phi=90°) at 17.7 GHz. The half-power beamwidth is 60°. The inset of Fig. 3 shows the simulated 3D radiation pattern of the optimized CDAA at 17.7 GHz. The gain in the direction of Oz-axis as a function of the frequency is depicted in Fig. 4. This gain is higher than 8.5 dBi in the analyzed frequency band (17 GHz -24 GHz) and reaches its maximum (11 dBi) at 22.3 GHz. The simulated input impedance viewed by the diode port of the CDAA (rectenna) is depicted in the Fig. 5. This impedance (Zin_a) should be the conjugate of the input impedance of the rectifier (Zin_r) in order to achieve the maximum power transfer between CDAA and the rectifier.

C. Design Methodology

A low-cost silicon Schottky diode (SMS201 from Aeroflex Metelics) in a molded plastic (DFN) package was adopted for this design. This diode is designed for a broadband zero bias detector. It can be used below 26.5 GHz for power detection up to 10 dBm. Most of previously reported designs below 10 GHz uses GaAs Schottky diodes with an higher cut-off frequency (e.g. M/A Com MA4E-1317 diode used in [1]- [START_REF] Takacs | Microwave Power Harvesting for Satellite Health Monitoring[END_REF], [START_REF] Ren | 35 GHz rectifying antenna for wireless power transmission[END_REF], [START_REF] Ladan | High Efficiency Low-Power Microwave Rectifier for Wireless Energy Harvesting[END_REF] has a cut-of frequency of 70 GHz while the Metelics MZBD-9161 diode used in [START_REF] Takacs | Microwave energy harvesting for satellite applications[END_REF]- [START_REF] Takacs | Compact Rectenna for Space Application[END_REF] has a cut-off frequency of 110 GHz). This low-cost silicon diode has also a good thermal behavior and can be mounted using a classical soldering process at high temperature (260°C) while the above-mentioned GaAs diodes require a more sensitive mounting process at a lower temperature (< 150°C).

The proposed design methodology combines full-wave electromagnetic simulation (for CDAA) and non-linear circuit simulation (for rectifier part). The main goal is to simultaneously achieve a high gain for the CDAA with an acceptable matching in a large bandwidth between CDAA and the rectifier for various load impedance and incident RF power. This was performed for maximizing the overall conversion efficiency of the rectenna. The design procedure consists of the following steps:

(i) the CDAA was first designed, simulated and optimized using FEKO software for operation in the desired frequency band. The diode was replaced by a voltage port while the capacitance (1.5 pF) and the load (e.g. 300 Ω) were modeled as port loads. The antenna radiation pattern and the distribution of the electric currents on the metallic strips were simulated in order to insure the proper radiation mechanism and to maximize the gain in the Oz-axis direction (that is, in the direction perpendicular to the rectenna surface);

(ii) the input impedance (Zin_a) as well as the gain along the Oz-axis (theta=0°; phi=0°) of the CDAA (simulation model shown in Fig. 1) was computed in the targeted frequency band as a function of the diode/capacitance mounting position;

non-linear (harmonic balance) simulation was performed by using the AWR software for the rectifier part of the rectenna. The coplanar stripline sections supporting the rectifier were modeled as a sub-circuit at the electromagnetic level because of the lack of an appropriate transmission line model in AWR. A generic non-physical diode model available in AWR was customized with the SMS201 diode parameters [10]. The input impedance (Zin_r) of the rectifier was determined as a function of the diode/capacitance mounting position;

(iv) the position the capacitance was adjusted to ensure the matching condition, that is Zin_a=Zin_r * where the asterisk denotes the complex conjugated.

III. RECTENNA: EXPERIMENTAL RESULTS

The rectenna was fabricated on Rogers 6002 substrate (relative permittivity: 2.94, loss tangent: 0.0012, thickness: 508 µm). The layout was accommodated with the manufacturing tolerances available in a university unit equipped for general (low frequency) PCB manufacturing. A metallic plane (size: ≈ 4 cm x 4 cm) was positioned at 1.2mm below the rectenna. A 1.2 mm thick Rohacell HF51 [START_REF] Popovic | Low-Power Far-Field Wireless Powering for Wireless Sensors[END_REF] (dielectric constant in the range of 1.08) layer was used as spacer between rectenna and its reflector. The main dimensions of the manufactured rectenna (see Fig. 1 and Fig. 2) are: lx=1.4mm, ly=1.8mm, lay=10mm, lax=8.3mm, lad=4.3mm, l=7 mm and lc=6 mm.

A. Experimental Setup

An experimental setup (shown in Fig. 6) was used in order to recreate the electromagnetic environment existing on antenna panels of the broadcasting satellites. A microwave signal generated from an Anritsu MG3694B generator was injected at the input of a horn antenna (VT220HA20-SK from Vector Telecom) which illuminated the rectenna under test with a linear polarized E-field. The power density (µW/cm 2 ) can be computed as a function of E-field effective value E (V/m) on the antenna surface or as a function of the RF power P t delivered to the transmitting horn antenna of gain G t and positioned at the distance d from the rectenna under test, as follows:

( 2 ) An automatic acquisition routine was implemented in Labview software from National Instruments to speed-up the acquisition process. The harvested DC voltage was measured with the microwave source OFF (DC noise: V_off) and with microwave source ON (DC voltage: V_on). The DC voltage of interest (V_DC) was then derived from the difference V_DC=V_on -V_off. The DC power can be computed from the measured DC voltage as long as the load is known. The measured loss due to the coaxial cable and connectors between antenna and the signal generator was in the range of 2.5 dB in the operating frequency band.

B. Simulation and Experimental Results

The frequency of 17.7 GHz was targeted by our design because available cartographic maps (for the expected E-field values on the K band antenna panels) indicated that an incident E-field up to 127 V/m may illuminate few regions of these panels [START_REF] Takacs | Microwave Power Harvesting for Satellite Health Monitoring[END_REF]. Nevertheless the adopted 'non-resonant' matching technique (between CDAA and the rectifier) can provide a wideband behavior to the rectenna. The simulated non-linear input impedance (obtained by using AWR software and harmonic balance technique) of the rectifier is represented in the Fig. 7. The estimated RF power of the input at f=17.7 GHz of rectifier by using Friis free-space formula [START_REF] Balanis | Antenna Theory: Analysis and Design, 3rd Edition[END_REF] (with injected power into the transmitting antenna: P t =23.5 dBm, measured gain of the transmitting horn antenna: G t ≈18dBi, receiving CDAA simulated gain antenna G r ≈9 dBi, distance between the input port of the transmitting and the receiving antenna d=25 cm) is in the range of 5 dBm. The simulated (AWR) nonlinear input impedance of the rectifier derived from Fig. 7 is Zin_r≈(76+42*j)Ω while the simulated (FEKO) input impedance of the CDAA (Fig. 5) is Zin_a≈(86-87*j) Ω. The measured DC power at the input of an (optimal) resistive load of 300 Ω by using the setup described in the section II. A (Fig. 6) is shown in the Fig. 8. These measurement results were obtained at f=17.7 GHz. The delivered power of the Anritsu MG3694B signal generator was P=25 dBm while the injected power on the coaxial input port of the transmitting horn antenna was in the range of 22.5 dBm. Due to the inherent limitations of the adopted design methodology and simulation models the obtained DC power at 17.7 is 0.52 mW. A maximum DC power of 0.9 mW on a 0.3kΩ load was obtained at f=23.15 GHz. At this frequency the gain of CDAA along the Oz-axis reaches a maximum value of 11 dBi while the simulated gain (along the Oz-axis) is 9 dBi at 17.7 GHz. The efficiency η (in %) of a rectenna can be computed by using the following worst-case definition [START_REF] Popovic | Low-Power Far-Field Wireless Powering for Wireless Sensors[END_REF]:

(3) where P DC is the harvested DC power, S is the incident electromagnetic power density and A G (in cm 2 ) denotes the area of the radiating surface. Here P t ≈22.5 dBm, G≈18 dBi and d=25 cm. Thus E≈73 V/m and S~1413 µW/cm 2 . By taken into account the overall area of the rectenna A G ≈2.5 cm 2 a worst-case conversion efficiency around of 15% at 17.7 GHz (harvested DC power: 0.9 mW) and 25 % at 23.15 GHz (harvested DC power: 0.9 mW) were obtained.

The proposed design methodology suffers from several limitations: (i) the impact of the nonlinear behavior of the diode on antenna is neglected during CDAA design as explained in section II.B; (ii) the impact on rectenna performances of the diode package and of the two via holes used to connect the diode (mounted on the top side of the rectenna) with the low-pass filtering section of the rectifier (located on the bottom side) is not taken into account in the AWR model; (iii) the simulation model adopted for the diode could be improved because key parameters of the diode (series resistance, ideality factor, etc.) are not provided in the datasheet [10]. Also, the PCB was fabricated in a university unit equipped for standard (low frequency) PCB low-cost manufacturing. The capacitance mounting position is critical and the exact mounting position is not well controlled by the adopted manual assembly process (tolerances in the range of 0.5 mm). Works are under run in order to overcome these limitations and to improve the global rectenna efficiency.

IV. CONCLUSION

An innovative rectenna topology and the adopted design methodology for K band application were presented. With only one low-cost silicon Schottky diode and using this ultracompact rectenna (2.5 cm 2 /1.5 square wavelengths at 23 GHz) a DC power up to 0.9 mW can be harvested for an incident Efield close to 73 V/m in the K band.

Fig. 1 .

 1 Fig. 1. Top view (left, not to scale) of the simulation model (FEKO) and a photo (top side) of the manufactured rectenna (right).

Fig. 2 .

 2 Fig. 2. Top view (left, not to scale) of the simulation model (PCB substrate is not represented) and bottom view (right, not to scale) of the rectenna.

Fig. 3 .

 3 Fig. 3. Simulated (FEKO) gain of the CDAA (xoz-plane: phi=0° and yozplane: phi=90°) at the frequency f=17.7 GHz.

Fig. 4 .

 4 Fig. 4. Simulated gain (FEKO) along Oz-axis versus frequency.

Fig. 5 .

 5 Fig. 5. Simulated input impedance (FEKO) of the CDAA (rectenna) structure viewed by diode port.

Fig. 6 .

 6 Fig. 6. Experimental setup used for characterizing the performances of manufactured rectennas .

Fig. 7 .

 7 Fig. 7. Simulated (AWR) nonlinear input impedance (traingle marker: real part; square marker: imaginary part) of the rectifier as function of the RF power at the rectifier input (frequency f=17.7 GHz).

Fig. 8 .

 8 Fig.8. Simulated (AWR) nonlinear input impedance (traingle marker: real part; square marker: imaginary part) of the rectifier as function of the RF power at the rectifier input (frequency f=17.7 GHz).

ACKNOWLEDGMENT

  This work was supported by CNES French Space Agency within the framework of R&T n° 115052 contracts. The electronic laboratory of University of Toulouse (Paul Sabatier) is acknowledged for PCB manufacturing. The authors acknowledge the support of the European Commission and of the National Funding Agency from France (Region Midi-Pyrénées) through the MNT-ERA.NET project MEMIS (MEMS based Millimeter wave Imaging System) for developing the simulation models of the Schottky diodes used in the design of our rectenna.


Christina Stabler

Fabrice Célarié

Tanguy Rouxel

René Limbach

Lothar Wondraczek

Ralf Riedel

Emanuel Ionescu
email: ionescu@materials.tu-darmstadt.de

Fabrice Celarie

Effect of

Keywords: Silicon oxycarbide, elastic properties, Poisson's ratio, plastic deformation, thermal annealing

come

Introduction

Amorphous silicon oxycarbides (SiOC) can basically be regarded as SiO2 glasses which are mechanically strengthened by the presence of tetrahedrally coordinated carbon atoms. Their synthesis relies on the thermal conversion of polyorganosiloxanes or precursors derived from sol-gel synthesis of organically modified alkoxysilanes [START_REF] Colombo | Polymer-derived ceramics: 40 years of research and innovation in advanced ceramics[END_REF][START_REF] Stabler | Silicon oxycarbide glasses and glass-ceramics: "All-Rounder" materials for advanced structural and functional applications[END_REF]. Via this procedure, SiOC glasses with varying chemical compositions are obtainable. Depending on the carbon content, a distinction is usually made between two groups of amorphous SiOC materials, i.e.,

A C C E P T E D M

A N U S C R I P T glasses with low amounts of segregated carbon and carbon-rich compositions (i.e. > 20 wt.% segregated carbon [START_REF] Colombo | Polymer-derived ceramics: 40 years of research and innovation in advanced ceramics[END_REF][START_REF] Blum | Synthesis and characterization of carbon-enriched silicon oxycarbides[END_REF][START_REF] Kleebe | SiOC ceramic with high excess free carbon[END_REF]). The chemical composition of SiOC glasses determines their structural and functional properties. Due to the remarkable high temperature stability, with a notable resistance against crystallization up to about 1500 °C [START_REF] Colombo | Polymer-derived ceramics: 40 years of research and innovation in advanced ceramics[END_REF], SiOC glasses have attracted a growing attention for applications in combustion engines or as coatings on turbine blades. In these, knowledge on the elastic properties and plasticity of SiOC glasses, like their Young's modulus, brittleness, hardness or creep behavior, is inevitable. In particular, this includes information on the mechanical performance under extreme conditions of high temperatures and extended exposure times.

SiOC glasses are X-ray amorphous and built-up of corner-sharing SiO4-xCx tetrahedral units (SiO4, SiO3C, SiO2C2, SiOC3 and SiC4) [START_REF] Brequel | Systematic Structural Characterization of the High-Temperature Behavior of Nearly Stoichiometric Silicon Oxycarbide Glasses[END_REF][START_REF] Widgeon | Navrotsky, 29Si and 13C Solid-State NMR Spectroscopic Study of Nanometer-Scale Structure and Mass Fractal Characteristics of Amorphous Polymer Derived Silicon Oxycarbide Ceramics[END_REF]. The abundance of each tetrahedral unit depends on the chemical composition and thus, the polymeric precursor used for the glass preparation. In addition to the covalently bonded C linked to Si (network carbon), typical SiOC glasses also contain a sp 2 -hybridized segregated carbon phase, which is homogeneously dispersed in the SiOC glass matrix [START_REF] Bois | Structural Characterization of Sol-Gel Derived Oxycarbide Glasses. 2. Study of the Thermal Stability of the Silicon Oxycarbide Phase[END_REF][START_REF] Rosenburg | High-Temperature Raman Spectroscopy of Nano-Crystalline Carbon in Silicon Oxycarbide[END_REF]. At temperatures above 1250 °C, phase separation of the amorphous SiOC glass starts, as evident from the vanishing signals for SiO3C, SiO2C2 and SiOC3 tetrahedral units in the 29 Si MAS NMR spectra [START_REF] Bois | Structural Characterization of Sol-Gel Derived Oxycarbide Glasses. 2. Study of the Thermal Stability of the Silicon Oxycarbide Phase[END_REF][START_REF] Soraru | Structural characterization and high-temperature behavior of silicon oxycarbide glasses prepared from sol-gel precursors containing Si-H bonds[END_REF]. This phase separation results in the formation of SiOC glass-ceramics within the system SiO2-SiC-C, where the residual glass matrix is composed of SiO2 with homogeneously dispersed β-SiC nanoparticles and segregated carbon. Controlling the polymeric precursor along with the synthesis conditions enables the preparation of SiOC materials with a tailored microstructure.

The partial substitution of two-fold coordinated O atoms by four-fold coordinated C atoms in SiOC glasses is known to improve the glass network connectivity and, by extension, to enhance the thermal and mechanical stability. In comparison to vitreous silica, SiOC glasses typically exhibit higher Young's modulus, hardness and glass transition temperature [START_REF] Soraru | Mechanical characterization of sol-gel-derived silicon oxycarbide glasses[END_REF][START_REF] Rouxel | High temperature behavior of a gel-derived SiOC glass: elasticity and viscosity[END_REF][START_REF] Rouxel | Creep viscosity and stress relaxation of gel-derived silicon oxycarbide glasses[END_REF][START_REF] Kleebe | Phase separation in an SiCO glass studied by transmission electron microscopy and electron energy-loss spectroscopy[END_REF][START_REF] Moysan | Mechanical characterization of a polysiloxane-derived SiOC glass[END_REF]. Similar trends have previously been reported also for SiOC glass-ceramics [START_REF] Renlund | Silicon oxycarbide glasses. Part II. Structure and properties[END_REF][START_REF] Papendorf | Hightemperature creep behavior of dense SiOC-based ceramic nanocomposites: microstructural and phase composition effects[END_REF][START_REF] Ionescu | High-temperature creep behavior of SiOC glass-ceramics: influence of network carbon versus segregated carbon[END_REF][START_REF] Stabler | High-temperature creep behavior of a SiOC glass ceramic free of segregated carbon[END_REF]. Although the creep rates and viscosity of SiOC glass-ceramics have been found as being determined by both the chemical composition and phase composition [START_REF] Ionescu | High-temperature creep behavior of SiOC glass-ceramics: influence of network carbon versus segregated carbon[END_REF][START_REF] Stabler | Influence of SiC/Silica and Carbon/Silica Interfaces on the High-Temperature Creep of Silicon Oxycarbide-Based Glass Ceramics: A Case Study[END_REF], information on the influence of the chemical composition and phase composition on the elastic properties of SiOC glass-ceramics remains limited. Moreover, only little is known about the impact of annealing on the elastic properties of SiOC glasses and glass-ceramics. Rouxel et al. [START_REF] Rouxel | High temperature behavior of a gel-derived SiOC glass: elasticity and viscosity[END_REF] studied the Young's modulus of a SiOC glass during a thermal treatment at temperatures up to 1400 °C. Upon the phase separation above 1250 °C, an irreversible increase of Young's modulus was observed, which has been attributed to the precipitation of β-SiC nanoparticles. Beside the technological significance of these findings, there is a fundamental interest on the underlying structural changes inside the glass network. Since the Young's modulus is intimately related to the atomic structure of glasses, i.e., strength and density of the interatomic bonds [START_REF] Makishima | Direct calculation of Young's modulus of glass[END_REF], it can serve as an indicator for structural modifications inside the glass network during annealing or even high-temperature creep experiments [START_REF] Rouxel | High temperature behavior of a gel-derived SiOC glass: elasticity and viscosity[END_REF].

Here, we report on the elastic properties and plasticity of a SiOC glass as well as a series of SiOC glass-ceramics with varying chemical and phase compositions. For selected SiOC

A C C E P T E D M A N U S C R I P T
glass-ceramics, the elastic properties were also monitored up to temperatures of around 1000 °C. This has also been done for SiOC glass-ceramics, which have been subjected to a prior thermal treatment comparable to previous high temperature creep experiments performed on these SiOC glass-ceramics [START_REF] Stabler | High-temperature creep behavior of a SiOC glass ceramic free of segregated carbon[END_REF][START_REF] Stabler | Influence of SiC/Silica and Carbon/Silica Interfaces on the High-Temperature Creep of Silicon Oxycarbide-Based Glass Ceramics: A Case Study[END_REF].

Experimental Procedure

Materials synthesis and processing. The synthesis procedure of the investigated SiOC glass and glass-ceramics is described in detail in Refs. [START_REF] Stabler | High-temperature creep behavior of a SiOC glass ceramic free of segregated carbon[END_REF] and [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF]. Different polymeric precursors were used to realize a series of SiOC materials with various amounts of segregated carbon. This includes one sample with no segregated carbon phase (denoted as SiC/SiO2), and further samples which contain approximately 1, 12 and 16 vol.% of segregated carbon, respectively (denoted as C1-SiOC, C12-SiOC, and C16-SiOC). SiC/SiO2 glass powder was prepared by pyrolysis of polysilsesquioxane in hydrogen atmosphere [START_REF] Narisawa | White Si-O-C(-H) particles with photoluminescence synthesized by decarbonization reaction on polymer precursor in a hydrogen atmosphere[END_REF]. C1-SiOC was prepared by a sol-gel synthesis of triethoxysilane and methyldiethoxysilane in a 2:1 molar ratio, followed by ageing and pyrolysis [START_REF] Soraru | Structural characterization and high-temperature behavior of silicon oxycarbide glasses prepared from sol-gel precursors containing Si-H bonds[END_REF]. C12-SiOC was synthesized via cross-linking and pyrolysis of a commercially available polysilsesquioxane (Belsil PMS MK, Wacker GmbH, Burghausen, Germany). C16-SiOC was prepared by a catalyzed sol-gel synthesis of a mixture consisting of 80 wt.% polymethylhydrosiloxane (PMHS; average Mn 1700-3200; Merck, Germany) and 20 wt.% 1,3,5,7-tetramethyl-1,3,5,7-tetravinylcyclotetrasiloxane (D4Vi; 97%, ABCr, Karlsruhe, Germany), followed by ageing and pyrolysis [START_REF] Radovanovic | Silicon oxycarbide glasses from silicone networks[END_REF]. The synthesized powders were ground and sieved (particle diameter < 40 µm) and were subsequently densified for 30 min at a temperature of 1600 °C and a pressure of 50 MPa using a uniaxial hot press operating in a static argon atmosphere (these samples were used for nanoindentation). Alternatively, the densification was achieved using a Field Assisted Sintering Technique (FAST). Here, the powders were treated at the same temperature and pressure but for only 15 min in argon atmosphere. The samples prepared by the latter routine were used in the Resonant Frequency Damping Analyzer (RFDA). The final monolithic samples are further denoted as SiC/SiO2-1600, C1-SiOC-1600, C12-SiOC-1600 and C16-SiOC-1600, respectively. Additionally, one monolithic SiOC glass (denoted as C1-SiOC-1100) was prepared from a sol-gel derived precursor of triethoxysilane and methyldiethoxysilane in a 2:1 molar ratio. The precursor was cast in plastic boxes, followed by curing and drying. The resulting monolithic xerogels were pyrolyzed in an alumina tube furnace for 3 h at 1100 °C in argon atmosphere (heating rate of 25 °C/h). Details on the preparation of C1-SiOC-1100 are reported in [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF].

Thermal treatment of the SiOC glass-ceramics were performed for different durations at temperatures of 1100, 1200 and 1300 °C, respectively, in an alumina oven with SiC heating elements using alumina crucibles with a closed lid under air. The applied heating rate was 20 °C/min up to 50 °C below the target temperature, followed by a slower rate of 2 °C/min up to the desired annealing temperature. The subsequent cooling was performed at a rate of 10 °C/min.

Structural characterization.

All samples were cut into smaller pieces using a diamond wire cutter and subsequently ground to co-planar geometry. Their chemical composition was

A C C E P T E D M A N U S C R I P T
determined using a carbon analyzer Leco-200 (Leco Corporation,USA) and a N/O analyzer Leco TC-436 (Leco Corporation, USA). The silicon elemental content was calculated as the difference to 100 wt.%, assuming no other elements being present in the samples. Archimedean (skeletal) density ρ and open porosity were derived upon water immersion technique. Powder X-ray diffraction (XRD) measurements were performed in flat-sample transmission geometry on a STOE STADI P diffractometer (Stoe, Darmstadt, Germany), equipped with a Mo X-ray tube and a position sensitive detector with a 6° aperture. Raman spectra were recorded on a Horiba HR800 micro-Raman spectrometer (Horiba JobinYvon, Bensheim, Germany), equipped with a He-Ne laser (λ = 633 nm). The measurements were conducted by using a grating of 600 g/mm and a confocal microscope (magnification 50 x NA0.75numerical aperture) with a 100 μm aperture, providing a resolution of approximately 1 μm. The laser power of 20 mW was attenuated by using neutral density filters.

Assessment of the elastic and plastic deformation behavior. Nanoindentation experiments were carried out on co-planar, optically polished samples using a G200 nanoindenter (Agilent Inc.), equipped with a three-sided Berkovich diamond indenter tip (Synton-MDP Inc.). Before the first experiments, both the tip area function and the instrument's frame compliance were calibrated on a Corning 7980 fused silica reference glass sample (Corning Inc.), following the procedure proposed by Oliver and Pharr [START_REF] Oliver | An improved technique for determining hardness and elastic modulus using load and displacement sensing indentation experiments[END_REF]. Depth profiles of the Young's modulus E and hardness H were obtained by operating in the continuous stiffness measurement mode [START_REF] Li | A review of nanoindentation continuous stiffness measurement technique and its applications[END_REF][START_REF] Pharr | Critical issues in making small-depth mechanical property measurements by nanoindentation with continuous stiffness measurement[END_REF]. For statistical relevance, at least ten indentations with a maximum displacement of 2 µm were created at a constant strain-rate of ε̇ = 0.05 s -1 (defined as the loading rate dP/dt divided by the actual load P). The strain-rate sensitivity m was analyzed via a nanoindentation strain-rate jump test as described in detail in Ref. [START_REF] Limbach | Strain-rate sensitivity of glasses[END_REF]. In total, ten strainrate jump tests with strain-rates of ε̇ = 0.05; 0.007 and 0.001 s -1 (in descending order) were performed. Values of m were derived from the slope of the linear regression between the logarithm of hardness and the logarithm of the indentation strain-rate (defined as the displacement rate dh/dt divided by the actual indentation depth h) m = ∂lnH/∂lnε̇𝑖 [START_REF] Shen | Nanoindentation creep of tin and aluminium: A comparative study between constant load and constant strain rate methods[END_REF], where 𝜀̇𝑖 = 𝜀̇2 ⁄ for materials with a depth-independent hardness [START_REF] Lucas | Indentation Power-Law Creep of High-Purity Indium[END_REF]. All nanoindentation experiments were carried out in laboratory air at ambient temperatures of around 301 ± 3 K and with thermal drift rates below 0.05 s -1 . To avoid interactions between residual stress fields, consecutive indentations were conducted at distances of 50 µm [START_REF] Hay | Introduction to Instrumented Indentation Testing[END_REF]. The same samples were subsequently utilized to study the elastic properties through ultrasonic echography with an Echometer 1077 (Karl Deutsch GmbH & Co. KG, Wuppertal, Germany). Values of the longitudinal cL and transversal sound wave velocities cT were calculated from the corresponding sound wave propagation times, as recorded with an accuracy of ± 1 ns by piezoelectric transducers operating at frequencies of 8 to 12 MHz, divided by the exact thickness of the glass plates, which was determined with an accuracy of ± 2 µm using a micrometer screw. On that basis, the shear modulus G, bulk modulus K, and Young's modulus, as well as the Poisson's ratio , were estimated according to the following equations [START_REF] Limbach | Plasticity, crack initiation and defect resistance in alkali-borosilicate glasses: From normal to anomalous behavior[END_REF]:

𝐺 = 𝜌𝑐 𝑇 2 (1) A C C E P T E D M A N U S C R I P T 𝐾 = 𝜌 (𝑐 𝐿 2 - 4 3 𝑐 𝑇 2 ) (2) 𝐸 = 𝜌 [ 3𝑐 𝐿 2 -4𝑐 𝑇 2 (𝑐 𝐿 /𝑐 𝑇 ) 2 -1 ] (3)
𝜈 = 𝑐 𝐿 2 -2𝑐 𝑇 2 2(𝑐 𝐿 2 -𝑐 𝑇 2 ) (4)
In addition, the temperature dependence of the Young's modulus and shear modulus were evaluated using a RFDA (IMCE NV RFDA HT 1050) on co-planar disk-like or rectangular bars with an approximate thickness of 2 mm. Samples were fixed with Pt-Rh wires and heated to 1000 °C at a rate of 10 °C/min in air. The subsequent cooling was carried out at a rate of 10 °C/h.

The investigation of the samples with both, RFDA and Ultrasonic Echography, requires sufficiently large specimens to be reliably fixed in the device or to detect the signal, respectively. The required minimum sample geometries could not be realized for samples C1-SiOC-1100 (due to the inherent challenges during the polymer-to-ceramic transformation) and SiC/SiO2-1600 (due to a limited amount of starting material). Consequently, the mentioned two methods were used only for samples meeting the respective requirements (see Table 1 below).

Results and Discussion

The as-prepared SiOC samples were first analyzed concerning their chemical and phase composition. For a detailed compositional and structural characterization the reader is referred to Refs. [START_REF] Stabler | High-temperature creep behavior of a SiOC glass ceramic free of segregated carbon[END_REF] and [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF]. The monolithic sample pyrolyzed at 1100 °C (i.e., C1-SiOC-1100) is fully X-ray amorphous and consists of an SiOC glass matrix with a homogeneously dispersed phase of sp 2 -hybridized segregated carbon [START_REF] Brequel | Systematic Structural Characterization of the High-Temperature Behavior of Nearly Stoichiometric Silicon Oxycarbide Glasses[END_REF]. The monolithic specimens prepared from hot pressing/FAST at 1600 °C can be regarded as SiOC glassceramics composed of a vitreous silica matrix with homogeneously dispersed β-SiC nanoparticles and segregated carbon [START_REF] Brequel | Systematic Structural Characterization of the High-Temperature Behavior of Nearly Stoichiometric Silicon Oxycarbide Glasses[END_REF][START_REF] Ionescu | High-temperature creep behavior of SiOC glass-ceramics: influence of network carbon versus segregated carbon[END_REF][START_REF] Stabler | High-temperature creep behavior of a SiOC glass ceramic free of segregated carbon[END_REF]. The phase composition in the system SiO2-SiC-C can be estimated from elemental analysis [START_REF] Widgeon | Navrotsky, 29Si and 13C Solid-State NMR Spectroscopic Study of Nanometer-Scale Structure and Mass Fractal Characteristics of Amorphous Polymer Derived Silicon Oxycarbide Ceramics[END_REF]. The chemical and phase compositions of the samples investigated in the present study are summarized in Table 1 and presented in Figure 1. The volume fractions of SiO2 and SiC for C1-SiOC-1100 can be regarded as the fraction of Si-O and Si-C bonds, respectively. The volume fraction of SiC is comparable in all samples. This is also evident from Figure 1, where the compositions investigated in this study lie on a parallel line to the SiO2-C tie-line in the Si-O-C ternary phase diagram. No segregated carbon was noticed in the SiC/SiO2-1600 sample. By comparing C1-SiOC-1100 with C1-SiOC-1600, it can be concluded that the phase separation has no influence on the overall chemical composition (e.g., through volatilization).

A C C E P T E D

M A N U S C R I P T Table 1.
Chemical compositions of the monolithic SiOC samples, their estimated phase compositions and the corresponding skeletal densities ρ. Volume fractions of SiO2, SiC and segregated carbon Csegr are calculated by means of the following densities: ρ(SiO2) = 2.20 g/cm 3 [START_REF] Renlund | Silicon oxycarbide glasses. Part II. Structure and properties[END_REF], ρ(β-SiC) = 3.22 g/cm 3 [START_REF] Bind | Phase transformation during hot-pressing of cubic silicon carbide[END_REF], ρ(Csegr) = 1.82 g/cm 3 (Graphitized Mesoporous Carbons GMC, Sigma-Aldrich, CAS-# 1333-86-4). The skeletal densities of the investigated SiOC samples are summarized in Table 1.

Replacing oxygen in the SiO2 network by carbon leads to an increase in density from 2.20 up to 2.28 g/cm 3 in the C1-SiOC-1100 glass specimen. The further increase in density from 2.28

A C C E P T E D M A N U S C R I P T
to 2.36 g/cm 3 is caused by the phase separation in the C1-SiOC-1600 sample and the accompanied precipitation of β-SiC nanoparticles (with a bulk density of 3.21 g/cm 3 for crystalline β-SiC [START_REF] Bind | Phase transformation during hot-pressing of cubic silicon carbide[END_REF]).

Along with the density, the Young's modulus is significantly enhanced upon the incorporation of carbon into vitreous silica (Table 2). The same trend has already been revealed in previous studies [START_REF] Soraru | Mechanical characterization of sol-gel-derived silicon oxycarbide glasses[END_REF][START_REF] Moysan | Mechanical characterization of a polysiloxane-derived SiOC glass[END_REF][START_REF] Walter | Microstructural and mechanical characterization of sol gel-derived Si-O-C glasses[END_REF], where it was attributed to the increased degree of cross-linking within the silica glass network [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF]. This effect even compensates the significantly lower energy of the Si-C interatomic bonds (U0 = 447 kJ/mol) compared to the strong Si-O interatomic bonds (U0 = 800 kJ/mol) in vitreous silica [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF]. Besides, phase separation does not seem to significantly alter Young's modulus. Actually, a change in Young's modulus during phase separation might be expected due to the reorganization of the network carbon in β-SiC nanoparticles and the accompanied formation of an oxygen-rich silica matrix. As a result, Young's modulus should decrease during phase separation, getting closer to the value of vitreous silica. However, Young's modulus of glass-ceramics is expected to be an additive function of the constituting phases and their respective volume fractions and properties [START_REF] Stryjak | Microstructure and properties of transparent glass-ceramics. Part 2. The physical properties of spinel transparent glass-ceramics[END_REF][START_REF] Miyata | Discussion of the indentation hardness of a glass-ceramic with participate microstructure[END_REF]. The comparably high Young's moduli of SiOC glass-ceramics are therefore a consequence of the high Young's modulus of the β-SiC nanoparticles (E = 300 GPa [START_REF] Soraru | High-Strength, High-Modulus Silicon-Oxycarbide glasses[END_REF]), compensating the smaller value of the vitreous silica matrix. The present data of the Young's modulus for the SiOC glass-ceramics are of similar magnitude as previously reported values for SiOC glasses [START_REF] Renlund | Silicon oxycarbide glasses. Part II. Structure and properties[END_REF][START_REF] Papendorf | Hightemperature creep behavior of dense SiOC-based ceramic nanocomposites: microstructural and phase composition effects[END_REF]. Figure 2a compares the available literature data for Young's modulus of SiOC materials at ambient temperature to the values determined in this study. Depending on the experimental technique used, Young's modulus of the SiOC samples with identical compositions varies within a range of ± 5 GPa (i.e., ± 5 %). Also, some of the literature values derived from

E [GPa] G [GPa] 𝜈 E [GPa] G [GPa] K [GPa] 𝜈 E [GPa] H [GPa] m C1-SiOC-1100 - - - - - - -
A C C E P T E D M A N U S C R I P T
bending or indentation experiments display relatively large standard deviations [START_REF] Soraru | Mechanical characterization of sol-gel-derived silicon oxycarbide glasses[END_REF][START_REF] Moysan | Mechanical characterization of a polysiloxane-derived SiOC glass[END_REF][START_REF] Walter | Microstructural and mechanical characterization of sol gel-derived Si-O-C glasses[END_REF]. However, a rough tendency is discernable in Figure 2a, that Young's modulus in SiOC glasses and glass-ceramics decreases with increasing content of sp 2 -hybridized segregated carbon. While the C1-SiOC-1100 glass sample with only 1 vol.% of sp 2 -hybridized segregated carbon investigated in this study exhibits a Young's modulus of 101.7 GPa, a much lower Young's modulus of only 66 GPa has previously been reported for a SiOC glass with approximately 64 vol.% of segregated carbon [START_REF] Martinez-Crespiera | Pressureless synthesis of fully dense and crack-free SiOC bulk ceramics via photo-crosslinking and pyrolysis of a polysiloxane[END_REF]. A plot of Young's modulus against the volume fraction of SiC in Figure 2b (for SiOC glasses, the volume fraction of SiC is thought to be systematically shifted to lower values due to the unknown density of the disordered arrangement of SiC4 tetrahedra) reveals their direct interrelation. SiOC materials with a larger fraction of SiC typically exhibit a higher Young's modulus as compared to samples containing a smaller fraction of SiC. This result is supposed to be related to the high values of E = 300 GPa of the nano-sized β-SiC precipitates [START_REF] Soraru | High-Strength, High-Modulus Silicon-Oxycarbide glasses[END_REF]. Thus, sp 2 -hybridized segregated carbon and β-SiC nanoparticles have the opposite effect on Young's modulus of SiOC glasses and glass-ceramics. This dependency was already suggested by Soraru et al. for a significantly smaller compositional range of SiOC glasses and for thin films [START_REF] Soraru | Mechanical characterization of sol-gel-derived silicon oxycarbide glasses[END_REF][START_REF] Sorarù | Influence of free carbon on the Young's modulus and hardness of polymer-derived silicon oxycarbide glasses[END_REF] and is extended now in this study to bulk SiOC glasses with higher amounts of segregated carbon and monolithic SiOC glass-ceramics.

Figure 3 illustrates the temperature dependence of Young's modulus of the investigated SiOC glass-ceramics with 1, 12 and 16 vol.% of segregated carbon, respectively. All samples are characterized by an increasing Young's modulus with increasing temperature. It is known from literature that Young's modulus of SiOC glasses and glass-ceramics increases with temperature (i.e. stiffening) analogous to vitreous silica [START_REF] Rouxel | High temperature behavior of a gel-derived SiOC glass: elasticity and viscosity[END_REF][START_REF] Papendorf | Hightemperature creep behavior of dense SiOC-based ceramic nanocomposites: microstructural and phase composition effects[END_REF][START_REF] Spinner | Temperature dependence of Young's modulus of vitreous germania and silica[END_REF]. In the case of vitreous silica, this behavior has been attributed to continuous atomic displacements during an amorphous-amorphous transformation comparable to the structural rearrangements associated to the α-to β-cristobalite phase transition in crystalline SiO2 [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF][START_REF] Lacks | First-Order Amorphous-Amorphous Transformation in Silica[END_REF]. We therefore conclude that the temperature dependence of the Young's modulus in the present SiOC glass-ceramics is dominated by the silica glass matrix. Since the Young's modulus did not show any hysteresis effects during cooling (not shown here), irreversible changes during the measurements can be excluded. Depending on the respective length scale, several different factors may affect the Young's modulus of glasses, including the atomic bonding energy and packing density, but also the network connectivity [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF] and super-structural heterogeneity [START_REF] Benzine | Local deformation of glasses is mediated by rigidity fluctuation and granularity[END_REF][START_REF] Sawamura | Scratch hardness of glass[END_REF]. Consequently, the stability of a glass, e.g., upon exposure to high temperatures can be monitored well by the measurement of Young's modulus, as a change in Young's modulus would be caused by a change in one or more of the above-mentioned structural parameters.

Thus, in the present study, Young's modulus was used to identify possible structural changes (i.e. decomposition or crystallization of the amorphous matrix) occurring in SiOC glassceramics (prepared at 1600 °C; i.e., fully phase-separated) which may result from their thermal annealing, for example during high-temperature creep experiments. Annealing temperatures and holding times were chosen according to the time necessary to reach the steady-state regime (secondary creep) in creep experiments, as previously reported [START_REF] Stabler | Influence of SiC/Silica and Carbon/Silica Interfaces on the High-Temperature Creep of Silicon Oxycarbide-Based Glass Ceramics: A Case Study[END_REF]. For sample C1-SiOC-1600, the chosen annealing conditions were 6 days at 1100 °C and 2 hours at 1300 °C, respectively, while for samples C12-SiOC-1600 and C16-SiOC-1600 the conditions were 6 days at 1200 °C and 12 hours at 1300 °C, respectively. Figures 4ac show the temperature evolution of Young's modulus before and after annealing at the chosen conditions. All SiOC glass-ceramics exhibit an increase in Young's modulus after annealing, where the annealing at higher temperature leads to a larger increase. In comparison to the as-prepared sample, C1-SiOC-1600 shows the lowest total increase of Young's modulus of about 1.4 %; whereas C12-SiOC-1600 exhibits an increase of about 7.3 %. C16-SiOC-1600 shows the highest total increase of the Young's modulus (approximately 8.8 % at 1200 °C). Interestingly, C16-SiOC-1600 shows lower Young's modulus (and its increase) during the thermal annealing at 1300 °C as compared to the features recorded during annealing at 1200 °C (see Figure 4c). The SiOC glass-ceramics are already fully phase-separated (i.e. the evolution of β-SiC nanoparticles occurred already during synthesis). Consequently, this increase in Young's modulus cannot be attributed to the precipitation of β-SiC nanoparticles.

In order to further rationalize the increase in Young's modulus and to directly compare all investigated samples, Figure 4d displays the temperature-dependent Young's modulus normalized to the corresponding value at room temperature of the investigated SiOC glassceramics (as-prepared and after thermal annealing). The temperature sensitivity of Young's modulus, i.e. the evolution of Young's modulus with temperature, is not affected by thermal annealing, as the curves of the individual glass-ceramics are identical within the estimated accuracy of the measurement (± 1 GPa; < 2 %). The temperature sensitivity can be correlated to the fragility of glasses [START_REF] Idriss | Evolution of the elastic modulus of Zr-Cu-Al BMGs during annealing treatment and crystallization: Role of Zr/Cu ratio[END_REF]. Furthermore, the fragility of glasses is connected to their atomic packing density, where glasses with high atomic packing density are known to be more fragile [START_REF] Angell | Perspective on the glass transition[END_REF]. Consequently, the stable temperature sensitivity indicates, that the glassy matrix of the investigated SiOC glass-ceramics is not significantly affected by thermal annealing. [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF] In general, two main strategies for improving Young's modulus for specific glass series are discussed in glass science: (i) an increase in the atomic packing density and (ii) an increase

A C C E P T E

A C C E P T E D M A N U S C R I P T
in the local bond strength [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF]. Both do not apply for our SiOC glass-ceramics. The first is expected to be related to an increase in the skeletal density and the second relies on a change in the elemental composition. For the investigated SiOC samples, the skeletal densities remain constant and consequently the first effect can be ruled out. Moreover, as the samples show only negligible weight loss and changes in their elemental composition (Table 3), also the second effect can be excluded. Consequently, the increase of Young's modulus if SiOC glass-ceramics after thermal annealing should have other origin and further supports the stability of the glassy matrix as already discussed considering the temperature sensitivity. One possible effect may arise from crystallization processes of the dispersed phases in SiOC glass-ceramics which take place during thermal annealing.

It is known that Young's moduli of nanocrystalline materials are significantly lower than those of their crystalline counterparts [START_REF] Soraru | High-Strength, High-Modulus Silicon-Oxycarbide glasses[END_REF]. Consequently, it is expected that an increased ordering /crystallization of the β-SiC nanoparticles or the segregated carbon phase in our SiOC samples might explain the observed increase in their Young's moduli. The evolution of the sp 2 -hybridized carbon phase in the SiOC glass-ceramics was studied by Raman spectroscopy, which may provide information on the average crystallite size La and the average distance between two defects LD in the segregated carbon phase. The procedure used for the determination of La and LD from the intensity ratio between D and G band (typical features in Raman spectra of disordered carbons) in silicon oxycarbides is described in detail in [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF] and values of the investigated samples are summarized in Table 4. Both La and LD are slightly increased in C1-SiOC-1600 and C16-SiOC-1600 after annealing at 1200 °C for 6 days; whereas, they were rather indifferent to the annealing experiment at 1300 °C for 12 hours. This indicates some increased ordering of the segregated carbon phase during the long-term annealing at 1200 °C, which may though explain the increase of the Young's modulus in SiOC after annealing. [START_REF] Stabler | Thermal Properties of SiOC Glasses and Glass Ceramics at Elevated Temperatures[END_REF] In addition to the sp 2 -hybridized carbon phase, the β-SiC nanoparticles may also be subjected to crystallization during the high-temperature annealing process. Figure 5 shows

A C C E P T E D M A N U S C R I P T
the XRD patterns of the SiOC glass-ceramics before and after thermal annealing. There is no indication for the crystallization of cristobalite. The presence of a reflection at 2θ = 11.5° for samples C16-SiOC-1600 annealed at 1200 °C for 6 days and C12-SiOC-1600 annealed at 1300 °C for 12 hours may be considered as an indication for the presence of the segregated carbon phase [START_REF] Mazo | Advanced silicon oxycarbide-carbon composites for high temperature resistant friction systems[END_REF]. This is in good agreement with the observed increase of La and LD as determined by Raman spectroscopy for C16-SiOC-1600. Differences between XRD and La and LD as determined from Raman spectroscopy for sample C12-SiOC-1600 are due to different volumes accessible for both measurement techniques, where powder XRD gives an averaged picture of the top layer of the whole sample, whereas Raman spectroscopy gives a localized information of low lateral size. Additionally, slight changes in the shape of the reflections related to β-SiC after thermal annealing can be identified. Rietveld refinement was performed with fixed lattice parameters. Crystallite size and micro-strain were refined according to the recommendations given in [START_REF] Balzar | Size-strain line-broadening analysis of the ceria round-robin sample[END_REF], with both Lorentzian and Gaussian contributions. For C1-SiOC-1600, a slight increase in average crystallite size of 2.2 to 3.1 nm and an increase in scale factor of about 20 % was determined after thermal annealing, indicating two effects caused by the high-temperature treatment of SiOC glass-ceramics. Firstly, an increase of the crystallite size of β-SiC reflects better ordering thereof. Secondly, an increase in scale factor relates to an increase of the

A C C E P T E D M A N U S C R I P T
amount of β-SiC, which in the present case can be explained either by the segregation of additional β-SiC during the thermal annealing (which would imply an incomplete phase separation of the SiOC glass during the hot-pressing step while preparing the monolithic glass-ceramics) or by the crystallization (i.e., increase in size) of precipitates that were too small to be reliably distinguished from the background of the XRD patterns (and are consequently not included in the scale factor for β-SiC). In other words, the increase of the scale factor is correlated to the crystallization of x-ray amorphous SiC to nanocrystalline β-SiC.

For C12-SiOC-1600 and C16-SiOC-1600, thermal annealing leads to a significant increase of the scale factor (ca. 200 % at maximum), accompanied by a slight decrease in the average crystallite sizes from 2.6-2.8 to 1.7-2.3 nm. The strong increase of the scale factor implies that there is a significant amount of X-ray amorphous SiC in the as-prepared monolithic SiOC glass-ceramics which crystallizes during the high-temperature annealing. At the same time, nanocrystalline β-SiC nanoparticles increase in size, too. Interestingly, as the amount of X-ray amorphous SiC nanoparticles which crystallizes is significant, there is an overall decrease of the average crystallite size after the thermal annealing.

From the above, it can be concluded that there is an increased ordering of β-SiC in all three samples upon thermal annealing (it is noted that the Rietveld refinement on the available data of such kind of disordered materials has to be regarded as being semi-quantitative and absolute values may be shifted). The most noticeable change is observed in C16-SiOC-1600 annealed at 1200 °C for 6 days. This is in line with the highest increase in Young's modulus and the highest carbon content. As it is suggested in literature that the presence of segregated carbon slows down β-SiC crystallization [START_REF] Brequel | Systematic Structural Characterization of the High-Temperature Behavior of Nearly Stoichiometric Silicon Oxycarbide Glasses[END_REF][START_REF] Brequel | Study of the phase separation in amorphous silicon oxycarbide glasses under heat treatment[END_REF], the as-prepared C16-SiOC-1600 sample is the least ordered sample with respect to β-SiC nanoparticles, and shows therefore the highest changes during thermal annealing. It is consequently suggested, that this increased ordering of the β-SiC nanoparticles is sufficient for an increase in Young's modulus, however, not high enough to be detectable at the scale of the skeletal density. As crystallization is a thermally activated process, higher temperatures and holding times favor higher degree of ordering and consequently Young's modulus. This can as well explain the relatively large increase in Young's modulus of C12-SiOC-1600 and C16-SiOC-1600 in comparison to C1-SiOC-1600. Thus, the effect of the β-SiC crystallization in SiOC-based glass-ceramics on their Young's modulus is significantly higher than that of their chemical composition.

A C C E P T E D

M A N U S C R I P T
The temperature dependencies of shear modulus, G, and Poisson's ratio, 𝜈, determined for C1-SiOC-1600 prior and post thermal annealing are shown in Figure 6 Figure .

In accordance to the evolution of Young's modulus, both shear modulus and Poisson's ratio are increasing with increasing temperature. This indicates that the mechanisms leading to a stiffening during temperature increase are at least partially resulting from an increase of the atomic packing density (as expressed by changes in Poisson's ratio), which is in line with the suggested structural rearrangements in vitreous silica [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF][START_REF] Lacks | First-Order Amorphous-Amorphous Transformation in Silica[END_REF].

Thermal annealing leads to a slight increase of the shear modulus, however within the estimated experimental error of ± 1 GPa. Poisson's ratio is not significantly impaired by thermal annealing as evidenced by comparing the values of 𝜈 before and after annealing at 1100 °C for 6 days. Poisson's ratio correlates to the atomic packing density, as will be addressed in the following, which in turn is expected to be mainly governed by the glassy matrix in a glass-ceramic. As Poisson's ratio of SiOC glass-ceramics is not significantly affected by the thermal treatment protocols, it is concluded that their silica-based amorphous matrix does not change during thermal annealing, indicating a good thermal stability of SiOC glass-ceramics. as determined by ultrasonic echography scatter only slightly within a narrow interval of 0.166 to 0.175, which compares very well with the range of Poisson ratios reported for vitreous silica of 0.15 to 0.18 [START_REF] Limbach | Strain-rate sensitivity of glasses[END_REF][START_REF] Brueckner | Mechanical properties of glasses[END_REF][START_REF] Pedone | Insight into Elastic Properties of Binary Alkali Silicate Glasses; Prediction and Interpretation through Atomistic Simulation Techniques[END_REF][START_REF] Zhao | Nanoindentation and Brillouin light scattering studies of elastic moduli of sodium silicate glasses[END_REF][START_REF] Sellappan | Composition dependence of indentation deformation and indentation cracking in glass[END_REF][START_REF] Sawamura | Lateral deformation and defect resistance of compacted silica glass: Quantification of the scratching hardness of brittle glasses[END_REF]. Interestingly, a SiOC glass derived from the same polymeric precursor as the C12-SiOC-1600 glass-ceramic was found to exhibit an extremely low Poisson's ratio of only 0.11 [START_REF] Moysan | Mechanical characterization of a polysiloxane-derived SiOC glass[END_REF]. Smaller Poisson's ratios in glasses are often correlated to strongly interconnected network structures of low atomic packing density [START_REF] Rouxel | Elastic properties and short-to medium-range order in glasses[END_REF][START_REF] Greaves | Poisson's ratio and modern materials[END_REF]. In case of SiOC glass, this is explained by the enhanced degree of cross-linking upon the exchange of two-fold coordinated oxygen atoms with four-fold coordinated carbon. The significantly higher Poisson's ratios of the SiOC glass-ceramics investigated in the present study are linked to the phase separation in SiOC glasses at temperatures exceeding 1250 °C [START_REF] Bois | Structural Characterization of Sol-Gel Derived Oxycarbide Glasses. 2. Study of the Thermal Stability of the Silicon Oxycarbide Phase[END_REF][START_REF] Soraru | Structural characterization and high-temperature behavior of silicon oxycarbide glasses prepared from sol-gel precursors containing Si-H bonds[END_REF]. We therefore conclude, that the Poisson ratio of our SiOC glass-ceramics are determined by the elastic properties of residual silica glass matrix.

The hardness of the studied SiOC glass and glass-ceramics scatter slightly around 10.5 to 11.4 GPa. The higher hardness of the C1-SiOC-1100 glass (H = 11.4 GPa) in comparison to vitreous silica (H = 9.3 GPa [START_REF] Limbach | Plasticity, crack initiation and defect resistance in alkali-borosilicate glasses: From normal to anomalous behavior[END_REF]) is the direct consequence of the enhanced network connectivity and the accompanied presence of additional network constraints, similarly to silicon oxynitride (SiON) glasses [START_REF] Paraschiv | Hardness of Oxynitride Glasses: Topological Origin[END_REF]. Before continuing this discussion, we need to note the mismatch in hardness obtained for the C1-SiOC-1100 glass (H = 11.4 GPa) investigated in the actual study and the previously published Vickers hardness values for SiOC glasses of 8.2 to 9.3 GPa [START_REF] Soraru | Mechanical characterization of sol-gel-derived silicon oxycarbide glasses[END_REF][START_REF] Kleebe | Phase separation in an SiCO glass studied by transmission electron microscopy and electron energy-loss spectroscopy[END_REF]. This discrepancy is attributed to the large contribution of elastic deformation to the indentation response of glasses. This manifests in marked differences between the contact area under load, which determines the hardness H, and the size of the residual hardness imprint after unloading, which is used for evaluating the Vickers hardness HV [START_REF] Oliver | Measurement of hardness and elastic modulus by instrumented indentation: advances in understanding and refinements to methodology[END_REF].

Unlike SiOC glasses, the improved hardness of the SiOC glass-ceramics is governed by a complex interplay between the mechanical stability, volume fraction and distribution of the residual silica glass matrix and the β-SiC nanoparticles as well as the segregated carbon [START_REF] Miyata | Discussion of the indentation hardness of a glass-ceramic with participate microstructure[END_REF]. Here, the β-SiC nanoparticles are supposed to be responsible for the improved hardness of the SiOC glass-ceramics, as polycrystalline β-SiC exhibits a Vickers hardness HV of 27.1 to 34.1 GPa [START_REF] Cha | Effects of activated reactive evaporation process parameters on the microhardness of polycrystalline silicon carbide thin films[END_REF]. Note, that the volume fraction of segregated carbon seems to have only little impact on the hardness. This is evident from the comparison of the SiC/SiO2-1600 (H = 10.6 GPa) and C12-SiOC-1600 (H = 10.5 GPa) glass-ceramics, which consist of an almost identical volume fraction of β-SiC nanoparticles, but zero and 12 vol.% segregated carbon, respectively. The same conclusion can be done upon comparing C1-SiOC-1600 (H = 11.0 GPa) and C16-SiOC-1600 (H = 10.6 GPa), revealing comparable amounts of β-SiC but 1 and 16 vol.% of segregated carbon, respectively. Interestingly, only marginal differences in hardness are distinguishable between the C1-SiOC-1100 glass (H = 11.4 GPa) and the C1-SiOC-1600 glass-ceramic (H = 11.0 GPa). This result is in clear contrast to a previous study, where an increasing hardness was noticed upon phase separation in a SiOC glass [START_REF] Kleebe | Phase separation in an SiCO glass studied by transmission electron microscopy and electron energy-loss spectroscopy[END_REF]. This was related to an increase in density, which is indeed occurring during phase separation (cf. Table 1), though effects related to the elimination of residual porosity in the SiOC glass were also considered. The brittleness of crystalline metals as well as metallic glasses can be correlated to the ratio between shear and bulk modulus (Pugh ratio G/K) [START_REF] Pugh | Relations between the elastic moduli and the plastic properties of polycrystalline pure metals[END_REF]. Metallic glasses show a tough-tobrittle transition at G/K > 0.41-0.43 [START_REF] Lewandowski | Intrinsic plasticity or brittleness of metallic glasses[END_REF]. In this context, Pugh ratios of 0.83 for C1-SiOC-1600, 0.86 for C12-SiOC-1600 and 0.85 for C16-SiOC-1600, would imply a brittle fracture behavior. The chemical composition has no significant influence on the Pugh ratio of SiOC glass-ceramics. Additionally, the index of brittleness B = H/KIc, which relates the resistance against plastic deformation to the resistance against fracture, was determined for C1-SiOC-1600 (B = 15 µm -0.5 ), C12-SiOC-1600 (B = 11 µm -0.5 ) and vitreous silica (B = 13 µm -0.5 ), where higher values correspond to more brittle material [START_REF] Lawn | Hardness, toughness, and brittleness: an indentation analysis[END_REF]. For the calculation the indentation hardness determined in this study was used, whereas values for fracture toughness KIc were taken from Ref. [START_REF] To | Fracture toughness and crack behavior of dense SiOC glass ceramics[END_REF]. As opposed to the Pugh ratio, the index of brittleness displays a notable dependence on the chemical composition. Sample C1-SiOC-1600 with only 1 vol.% of segregated carbon appears to be more brittle than vitreous silica. However, the presence of 12 vol.% segregated carbon in C12-SiOC-1600 has the opposite effect and drastically reduces the brittleness even below the level of vitreous silica. This behavior is in line with the experiences made during grinding of the samples and clearly signifies the importance of both the chemical and phase composition for the mechanical performance of SiOC glasses and glass-ceramics. This argumentation is also supported by previous observations regarding the effects of phase separation on the mechanical stability of alkali-alkaline earth silicate or alkali-borosilicate glasses [START_REF] Limbach | Plasticity, crack initiation and defect resistance in alkali-borosilicate glasses: From normal to anomalous behavior[END_REF][START_REF] Seal | Effect of phase separation on the fracture toughness of SiO2-B2O3-Na2O glass[END_REF][START_REF] Deriano | Structure and mechanical properties of alkalialkaline earth-silicate glasses[END_REF][START_REF] Cheng | Indentation cracking behaviour and structures of nanophase separation of glasses[END_REF]. Tuning the volume fraction and structure of the separated phases is therefore supposed as a promising route towards the fabrication of SiOC materials with an enhanced toughness.

For a potential application of SiOC glasses and glass-ceramics at elevated temperatures, a high creep resistance, i.e., low strain-rate sensitivity, is of paramount importance. Using a nanoindentation strain-rate jump test [START_REF] Limbach | Strain-rate sensitivity of glasses[END_REF], very low strain-rate sensitivities of 0.0059 to 0.0107 were determined for the SiOC glass and glass-ceramics. The results obtained for both the C1-SiOC-1100 glass (m = 0.0059) and the SiO2/SiC-1600 (m = 0.0060) glassceramic are very close to the values reported for vitreous silica (m = 0.0068 -0.0150 [START_REF] Limbach | Strain-rate sensitivity of glasses[END_REF][START_REF] Peykov | Evaluation of strain rate sensitivity by constant load nanoindentation[END_REF][START_REF] Elmustafa | Strain rate sensitivity in nanoindentation creep of hard materials[END_REF]), which displays no significant creep deformation at room temperature [START_REF] Chudoba | Investigation of creep behaviour under load during indentation experiments and its influence on hardness and modulus results[END_REF][START_REF] Fischer-Cripps | A simple phenomenological approach to nanoindentation creep[END_REF]. With respect to this result, we speculate the creep response of SiOC materials with low amounts of segregated carbon to be determined by the creep resistance of the vitreous silica glass matrix. Nevertheless, regarding the SiOC glass-ceramics a slight but still detectable increase of the strain-rate sensitivity from 0.0060 (SiO2/SiC-1600) to 0.0107 (C12-SiOC-1600) is seen with increasing amounts of segregated carbon. A higher strain-rate sensitivity in glasses normally signifies a homogenization of the plastic flow [START_REF] Meinhard | Hardness and flow behavior of glass in the nanometer range. An interpretation of the load dependence of the hardness[END_REF]. With the presence of a homogeneously dispersed phase of segregated carbon in the vitreous silica glass matrix, fertile sites for a shear-mediated plastic flow are created inside the material.

Conclusions

Silicon oxycarbides are known to possess increased mechanical properties like Young's modulus and hardness. The present study rationalizes the various impacts on their elastic 1. Chemical compositions of the monolithic SiOC samples, their estimated phase compositions and the corresponding skeletal densities ρ. Volume fractions of SiO2, SiC and segregated carbon Csegr are calculated by means of the following densities: ρ(SiO2) = 2.20 g/cm 3 [START_REF] Renlund | Silicon oxycarbide glasses. Part II. Structure and properties[END_REF], ρ(β-SiC) = 3.22 g/cm 3 [START_REF] Bind | Phase transformation during hot-pressing of cubic silicon carbide[END_REF], ρ(Csegr) = 1.82 g/cm 3 (Graphitized Mesoporous Carbons GMC, Sigma-Aldrich, CAS-# 1333-86-4). 3. Skeletal density ρ and elemental composition of the investigated as-prepared SiOC glass-ceramics and after a thermal treatment. Table 4. Effects of a thermal treatment on the degree of graphitization of the segregated carbon in SiOC glass-ceramics as determined from Raman spectroscopy: Average crystallite size La, average distance between two defects LD. AD/AG describes the ratio between the areas assigned to the D and G band, respectively.

Figure 1 .

 1 Figure 1. Ternary phase diagram in the system Si-O-C. The dashed line (parallel to the SiO2-C tieline) marks that the segregated carbon content increases from SiC/SiO2-1600 to C16-SiOC-1600 without large variations in the SiC content.

Figure 2 .

 2 Figure 2. Room temperature Young's modulus of SiOC glasses (open symbols) and glassceramics (filled symbols) as a function of (a) volume fraction of sp 2 -hybridized segregated carbon and (b) the volume fraction of Si-C bonds (as for SiOC glasses) or ß-SiC (as for SiOC glass-ceramics). Literature values (black symbols) are added for comparison. Dashed lines were drawn as a guide for the eyes.

Figure 3 .

 3 Figure 3.Temperature dependence of the Young's modulus of the SiOC glass-ceramics investigated in this study. The experimental error is estimated to be ± 1 GPa (i.e., < 2 %).

Figure 4 .

 4 Figure 4. Temperature dependence of the Young's modulus of the as-prepared (a) C1-SiOC-1600, (b) C12-SiOC-1600 and (c) C16-SiOC-1600 SiOC glass-ceramics, respectively, and after a thermal treatment at different temperatures and for varying durations. Lines drawn are fits to the experimental data. (d) Temperature dependence of the Young's modulus normalized to its value at room temperature (displayed by trendlines). The experimental error is estimated to be ± 1 GPa (i.e., < 2 %).

Figure 5 .

 5 Figure 5. X-ray diffraction patterns (Mo Kα) of SiOC glass-ceramics prior and after annealing at high temperatures: (a) C1-SiOC-1600; (b) C12-SiOC-1600; (c) C16-SiOC-1600. Indexed reflections relate to β-SiC.

Figure 6 .

 6 Figure 6.

  and plastic properties caused by the chemical/phase compositional variability inherent to the SiOC system and gives further insight in the thermal stability of SiOC glass-ceramics. The elastic properties of silicon oxycarbides are influenced by both phase composition (i.e. volume fraction of β-SiC or fraction of Si-C bonds in Si-O-C glass network and volume fraction of segregated carbon) as well as by their microstructure (i.e. SiOC glass vs. glassceramic; phase separation). Whereas Young's modulus is increasing with increasing volume fraction of β-SiC nanoparticles/Si-C bonds and decreasing with increasing amount of segregated carbon, it is not significantly affected upon phase separation. In contrast, Poisson's ratio, and consequently the atomic packing density, is independent of the chemical composition, but is significantly influenced by the phase separation. Both, the development of Young's modulus and Poisson's ratio, can be considered as direct consequences of: (i) the highly cross-linked network architecture of SiOC glasses due to the presence of mixed-bonds SiO4-xCx tetrahedra; (ii) the decrease of the cross-linking degree of the glassy network during the phase separation and (iii) the high stiffness of crystalline β-SiC. Thermal annealing of SiOC glass-ceramics leads to an increase in Young's modulus. It is demonstrated that this relies in an increased ordering of the dispersed phases segregated carbon and β-SiC as evidenced by XRD and Raman spectroscopy and not in a change of the glassy silica matrix as testified by an unbiased Poisson's ratio and temperature sensitivity of Young's modulus. This implies, that the silica-rich glassy matrix of SiOC glass-ceramics is stable at high temperatures and longer holding times. The plastic properties of SiOC glass-ceramics are significantly influenced by the volume fractions of both segregated carbon and β-SiC nanoparticles. The segregated carbon phase reduces brittleness and induces additional sites for sliding during plastic deformation; whereas, the hardness is governed by the volume fraction of β-SiC nanoparticles due to their comparably high hardness. Interestingly, the hardness in SiOC seems to be not affected by phase separation.

Figure 1 .

 1 Figure 1. Ternary phase diagram in the system Si-O-C. The dashed line (parallel to the SiO2-C tie-line) marks that the segregated carbon content increases from SiC/SiO2-1600 to C16-SiOC-1600 without large variations in the SiC content.

Figure 2 .Figure 3 .

 23 Figure 2. Room temperature Young's modulus of SiOC glasses (open symbols) and glassceramics (filled symbols) as a function of (a) volume fraction of sp 2 -hybridized segregated carbon and (b) the volume fraction of Si-C bonds (as for SiOC glasses) or ß-SiC (as for SiOC glass-ceramics). Literature values (black symbols) are added for comparison. Dashed lines were drawn as a guide for the eyes. Figure 3. Temperature dependence of the Young's modulus of the SiOC glass-ceramicsinvestigated in this study. The experimental error is estimated to be ± 1 GPa (i.e., < 2 %).

Figure 4 .

 4 Figure 4. Temperature dependence of the Young's modulus of the as-prepared (a) C1-SiOC-1600, (b) C12-SiOC-1600 and (c) C16-SiOC-1600 SiOC glass-ceramics, respectively, and after a thermal treatment at different temperatures and for varying durations. Lines drawn are fits to the experimental data. (d) Temperature dependence of the Young's modulus normalized to its value at room temperature (displayed by trendlines). The experimental error is estimated to be ± 1 GPa (i.e., < 2 %).Figure 5. X-ray diffraction patterns (Mo Kα) of SiOC glass-ceramics prior and after annealing at high temperatures: (a) C1-SiOC-1600; (b) C12-SiOC-1600; (c) C16-SiOC-1600. Indexed reflections relate to β-SiC. Figure 6. Temperature dependence of the shear modulus G and Poisson's ratio 𝜈 of the (a) as-prepared C1-SiOC-1600 SiOC glass-ceramic and (b) after a thermal treatment for 6 days at 1100 °C.

Figure 5 .Figure 6 .

 56 Figure 4. Temperature dependence of the Young's modulus of the as-prepared (a) C1-SiOC-1600, (b) C12-SiOC-1600 and (c) C16-SiOC-1600 SiOC glass-ceramics, respectively, and after a thermal treatment at different temperatures and for varying durations. Lines drawn are fits to the experimental data. (d) Temperature dependence of the Young's modulus normalized to its value at room temperature (displayed by trendlines). The experimental error is estimated to be ± 1 GPa (i.e., < 2 %).Figure 5. X-ray diffraction patterns (Mo Kα) of SiOC glass-ceramics prior and after annealing at high temperatures: (a) C1-SiOC-1600; (b) C12-SiOC-1600; (c) C16-SiOC-1600. Indexed reflections relate to β-SiC. Figure 6. Temperature dependence of the shear modulus G and Poisson's ratio 𝜈 of the (a) as-prepared C1-SiOC-1600 SiOC glass-ceramic and (b) after a thermal treatment for 6 days at 1100 °C. A C C E P T E D M A N U S C R I P T

Table 2 .

 2 Room temperature elastic properties and plasticity of SiOC glasses and glass-

	Sample	RFDA	Ultrasonic Echography	Nanoindentation

ceramics: Young's modulus E, shear modulus G, bulk modulus K, Poisson's ratio 𝜐, hardness H and strain-rate sensitivity m. Literature data for a SiOC glass and Suprasil vitreous silica are added for comparison.

Table 3 .

 3 Skeletal

	Sample	Annealing Conditions	Density ρ [g/cm 3 ]	Weight loss [%]	Elemental Composition Si [wt.%] O [wt.%] C [wt.%]
		as-prepared	2.38*	-	51.78*	41.59*	6.63*
	C1-SiOC-1600	1100 °C, 6 d	2.38	1.66	52.66	40.57	6.77
		1300 °C, 2 h	2.38	-	52.54	40.59	6.87
		as-prepared	2.31*	-	46.36*	39.52*	14.12*
	C12-SiOC-1600	1200 °C, 6 d	2.32	0	47.48	38.97	13.55
		1300 °C, 12 h	2.32	0	46.81	39.32	13.87
		as-prepared	2.34*	-	46.79*	33.82*	19.39*
	C16-SiOC-1600	1200 °C, 6 d	2.37	0.14	47.35	33.56	19.09
		1300 °C, 12 h	2.38	0.19	46.89	33.66	19.45
	* Values taken from Ref.

density ρ and elemental composition of the investigated as-prepared SiOC glassceramics and after a thermal treatment.

Table 4 .

 4 Effects of a thermal treatment on the degree of graphitization of the segregated carbon in SiOC glass-ceramics as determined from Raman spectroscopy: Average crystallite size La, average distance between two defects LD. AD/AG describes the ratio between the areas assigned to the D and G band, respectively.

	Sample	Annealing conditions	AD/AG	La [nm]	LD [nm]
		as prepared	3.185 ± 0.233	12.1 ± 0.9	9.5 ± 0.4
	C1-SiOC-1600	1100 °C, 6 d	2.800 ± 0.142	13.8 ± 0.7	10.2 ± 0.3
		1300 °C, 12 h	2.232 ± 0.356	17.3 ± 2.7	11.4 ± 0.9
		as prepared	4.215 ± 0.251*	9.2 ± 0.6*	8.3 ± 0.3*
	C12-SiOC-1600	1200 °C, 6 d	4.539 ± 0.422	8.5 ± 0.8	8.0 ± 0.4
		1300 °C, 12 h	4.699 ± 0.036	8.2 ± 0.1	7.8 ± 0.1
		as prepared	4.121 ± 0.521*	9.5 ± 1.3*	8.4 ± 0.6*
	C16-SiOC-1600	1200 °C, 6 d	3.093 ± 0.203	12.5 ± 0.8	9.7 ± 0.3
		1300 °C, 12 h	4.104 ± 0.075	9.4 ± 0.2	8.4 ± 0.1

* Values are taken from Ref.

Table Captions Table

 Captions

Table 2 .

 2 Room temperature elastic properties and plasticity of SiOC glasses and glassceramics: Young's modulus E, shear modulus G, bulk modulus K, Poisson's ratio 𝜐, hardness H and strain-rate sensitivity m. Literature data for a SiOC glass and Suprasil vitreous silica are added for comparison. Table

Molar fractions of SiO2 and SiC can be regarded as the fraction of Si-O and Si-C bonds, respectively.

Volume fractions of SiO2 and SiC can be regarded as the fraction of Si-O and Si-C bonds, respectively. Absolute values are expected to be shifted due to unknown density of the disordered arrangement of SiC4 tetrahedra.

Acknowledgements

The authors are grateful for financial support provided from the Deutsche Forschungsgemeinschaft (IO 64/7-1, "High-Temperature Creep in SiOC-Based Glasses and Glass-Ceramics").


A Takacs
email: takacs@laas.fr

D Dragomirescu

S Charlot
email: charlot@laas.fr

P-F Calmon
email: calmon@laas.fr

Flexible technology for millimeter-wave wireless sensors applications

Keywords:

This paper addresses recent advances obtained in LAAS CNRS concerning the developments of the flexible kapton technology and the use of this technology in the manufacturing of flexible and conformal structures for applications like Wireless Sensor Networks (WSNs) and Structure Health Monitoring (SHM). Flexible kapton supported passive devices are presented as well as the technological process developed for the heterogeneous integration of active circuits with such flexible passive devices.

I. INTRODUCTION

The developments of any successful complex industrial product in the field of aeronautics, space, automotive or IoT industries requires the use of a huge numbers of sensors deployed both on the final product or on the prototypes used during the development cycle. These sensors should be lowcost, low-weight and deployed on various surfaces. In order to deploy a large number of such sensors in a conformal manner on irregular surface a flexible substrate is required. Many flexible substrate technologies are today available [START_REF] Wong | Flexible Electronics: Materials and Applications[END_REF]- [START_REF]Toray Research Center, Printing Technology for Flexible Substrates[END_REF] and we focused our research on the kapton flexible substrates that offers a best trade-off between mechanical, thermal and electrical properties [START_REF] Dupont | DuPont TM Kapton ®[END_REF].

II. FLEXIBLE KAPTON TECHNOLOGY

To manufacture our circuits onto a flexible substrate, we have chosen the traditional photolithography mainly because of the intrinsic high resolution [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF]. The main difficulty during the fabrication process lies on the flexible film handling and its use in micro-technology equipment. In order to overcome this difficulty, a 4-inch silicon wafer is used as a host carrier. One of the critical obstacles consists of finding a way to adhere the polyimide film on the Si support. This adhesion has to be compatible with the various technological stages (vacuum, solvent and temperature) and allows after manufacturing a peeling without any physical or chemical constraint. A matured fabrication process was used and a PDMS (PolyDiMethylSiloxane) spin coating was performed for the adhesion of the polyimide on the holding wafer. Then the Kapton polyimide was patterned on the PDMS-Si support using a Shipley 360N laminator. A resin spin coating was then realized in a fully automated resist processing tool (EVG120), followed by a photolithography process. During the different fabrication processes, Kapton polyimide was metallized with Ti/Au layers (50 nm/ 200 nm) or Ti/Cu layers (50 nm/500 nm) using the Electron Beam Physical Vapor Deposition (EBPVD). In case of Cu coating, the fabrication process was terminated with a surface finishing (gold immersion deposits) to prevent Cu from oxidizing. Both wet-etching (Fig. 1) and lift-off methods (Fig. 2) can be used [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF], each of them allowing sufficient resolution (the minimum resolution is in the range of 1µm both the metallic strips and slot) for the metallization tracks for millimeter wave applications (up to W-band). The quality (homogeneity and roughness) of the metallization layers obtained by the adopted photolithography based technology are also better as compared with other flexible technology (e.g. inkjet printing). (1)

III. DESIGN OF PASSIVE KAPTON COMPONETS FOR MILLIMETER WAVE APPLICATIONS

In order to validate our technological approach a lot of passive components were designed, fabricated and manufactured in our laboratory [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF]- [START_REF] Yang | Design of Kapton based passive circuits at microwave frequencies[END_REF]. A thickness of 125 µm was selected for the Kapton film as a best trade-off between: (i) the commercially available thickness of Kapton films, (ii) the electromagnetic performances required by typical antennas (e.g. patch antennas) that can be used for millimeter wave applications, (iii) the mechanical properties. Fig. 3 shows a 4 inch kapton wafer including: (i) circular ring resonators (CRR), (ii) cross dipole antennas (by using strip or slot based topologies), (iii) patch antennas, (iv) rectennas (an acronym from rectifier antenna). CRR are useful structures for the characterization of Kapton dielectric properties (measurement of the relative electric permittivity ε r and dielectric loss tangent). The measurement results obtained with the CRR (designed to cover the frequency band up to 67 GHz) represented in Fig. 3 allows us to extract the dielectric constants of the kapton substrate. The relative dielectric permittivity is ε r = 3,2 ± 0.03 and the loss tangent is tanδ= 0.012 ± 0.004.

The strip (printed) cross dipoles antenna array (CDAA) exhibits good performances in terms of input matching and radiation pattern for RF application (e.g. WSN operating in the ISM 24 GHz band or energy harvesting devices/rectennas for space applications [START_REF] Takacs | Microwave power harvesting for satellite health monitoring[END_REF]) while the patch antennas and the cross slot dipole antennas (CSDA) were designed for V-band and Wband applications. Very interesting results were obtained with CDA and CSDA both in terms of frequency bandwidth and radiation pattern. CDA consists of two crossed dipoles printed on the kapton surface. For example the CDAA represented in Fig. 4 consist of an array formed by two cross dipole antennas. The two crossed dipoles are connected together by using coplanar striplines (CPS). The CDAA is equipped with an endlaunch coaxial (K) connector from Southwest. A special designed CPS-to-microstrip T junction is used to accommodate the CPS lines to the end-launch connector. The measured frequency bandwidth (S11<-10dB) is approximately 4 GHz (Fig. 5) and this antenna can be used as wide-bandwidth antenna in the K-band (e.g. energy harvesting of the spill-over losses of microwave antennas onboard of geostationary broadcasting satellites [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF]) or in the 24 GHz ISM band. The radiation pattern do not change significantly in the relative bandwidth (approximately 18% for K-band/ISM 24 GHz CDAA) and the measured maximum gain (Fig. 6) is about 4.3 dBi (including the losses added by the end lunch connector, the associated mounting process and by CPS-to-microstrip T junction). We note that the gain can be increased if a metallic plate acting as reflector is properly putted below the CDAA. A cross slot dipole antenna (CSDA) was also designed for V-band application. As represented in Fig. 7, CSDA can be obtained from the previously presented printed (strip) CDA topology by applying the Babinet principle. Instead of the radiative strip patterned on the Kapton surface we have now radiative slots. CDSA antenna is feed by using a coplanar waveguide line with a characteristic impedance of 50Ω at 60 GHz. The center frequency of the V-band CSDA is 60.2 GHz, and the -10 dB bandwidth is from 58.2 GHz to 62.05 GHz (6% of relative bandwidth). A peak antenna gain of 4 dB was obtained at 60.2 GHz. A parametric study for the angle (α) between slots was also performed. When the angle α increases, the peak antenna gain decreases and the half power beamwidth (HPBW) increases. This phenomenon can be explained by the increase of the coupling effect when the radiative slots approach each other.

Another key issue for any flexible technology is the technological solution/process to heterogeneous integrate active circuits on such flexible passive structures (e.g. to integrate a CMOS transmitter with an antenna and sensors in order to implement a wireless sensor node for SHM applications). A 3D heterogeneous approach based on a flipchip process was adopted in our laboratory [START_REF] Jatlaoui | 3D Heterogeneous Integration of Wireless Communicating Nano-Sensors on Flexible Substrate[END_REF]. In order to implement the flip-chip process two approaches are possible: (i) a flip-chip with Au stud bumps implemented on the pads of the integrated circuit (chip), (ii) a flip-chip process with Au electroplated bumps implemented on the kapton surface. The first approach (by using Au stud bumps) was initially developed in our laboratory [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF], [START_REF] Jatlaoui | 3D Heterogeneous Integration of Wireless Communicating Nano-Sensors on Flexible Substrate[END_REF] but have a major drawback: the gold stud bumping is a semi-manual process that lacks of reproducibility and reliability mainly when a large number of gold stud bumps/chip is required. The second method (with Au bumps developed on the Kapton surface) leads to a more reliable approach and is under development in our laboratory [START_REF] Yang | Flexible Substrate Technology for Millimeter Wave Applications[END_REF]. The Au bumps are created by using a technological process with the mains steps presented in Fig. 8. As shown in Fig. 9 a), we can assure a very good quality and alignment for the gold bumps directly fabricated/patterned on the Kapton. The Kapton structure represented in Fig. 9 a) was designed in order to accommodate a CMOS V-band low noise amplifier (LNA) designed in our laboratory and fabricated by using the 65 nm CMOS process from ST Microelectronics. A lot of tests were performed by using dummy chip fabricated on glass substrate. The dummy chips have a size (400 µm x 400 µm) and a pad configuration identically with the 'real' LNA. A non-conductive adhesive polymer is deposited on the inner surface delimited by the Au bumps and the chip is stick by thermo-compression on the Kapton surface. The main challenge is how to deposit a non-conductive adhesive polymer on a surface as small as 240 µm × 250 µm on Kapton. Unfortunately up to now this method do not lead to 100% reliable process. During our multiple test, at least one contact (between the gold bump and the pads of the dummy chip) was negatively impacted by the presence of the nonconductive polymer. Works are underrun in order to find a more reliable procedure. An alternative method to overcome the adhesive issue can be the use of a raised-die under-fill configuration for the flip-chip [START_REF] Zhang | Recent Advances in Flip-Chip Underfill: Materials Process and Reliability[END_REF].

IV. CONCLUSION

A kapton flexible technological process developed in our laboratory was presented. This process allow to obtain a very good resolution and quality for the metalization layers. Millimeter wave passive devices was manufactured by using this process. Very good performances were obtained for Kband cross dipole antenna array and V-band cross slot dipole antenna. Two technological process, under development in our laboratory for a flip-chip heterogenous integration technique that take into account the worst case constraints for millimeter wave applications, were also presented.

Fig. 1 .

 1 Fig. 1. Main technological steps of the wet-etching method

Fig. 3 .Fig. 4 .Fig. 5 .Fig. 6 .

 3456 Fig. 3. Manufactured Kapton wafer with double sided metallic patterns

Fig. 7 .

 7 Fig. 7. Top view (not to scale) of the V-band CSDA and the simulated (HFSS) 3D radiation patern

Fig. 8 .

 8 Fig. 8. Technological steps when the gold bumps are fabricated directly on the Kapton surface

Fig. 9 .

 9 Fig. 9. a) Photo of the gold bumps patterned on the kapton surface; b) photo of the dummy chip mounted by a flip-chip process on the kapton surface


M Gastaldi

D Dragomirescu

A Takacs

V Armengaud

S Rochette

Micro-strip slow wave line for phase shifting cells

This paper addresses the design, simulation and measurements of microstrip slow wave lines implemented in a 0.25 µm SiGe BiCMOS process for Ku-band (10 GHz -15 GHz) applications. The simulation results and the measurements show better performances for the proposed microstrip slow wave line compared to the classical microstrip transmission line. These lines present very low insertion losses and a high phase constant.

Introduction: Slow Wave Lines (SWL) are a specific topology of transmission lines that uses the slow wave phenomenon in order to miniaturize the physical length of a line [START_REF] Cheung | Shielded passive devices for siliconbased monolithic microwave and millimeter-wave integrated circuits[END_REF][START_REF] Tang; Franc | Performance Improvement Versus CPW and Loss Distribution Analysis of Slow-Wave CPW in 65 nm HR-SOI CMOS Technology[END_REF]. Nowadays, CMOS and BiCMOS (Bipolar CMOS) technologies are more and more used for microwave circuits [START_REF] Heydari | Low-Power mm-Wave Components up to 104GHz in 90nm CMOS[END_REF] and become a promising candidate for core-chip circuits used in the satellite Beamforming Antenna Control Systems (BACS). Mainly for the phase shifter of BACS the use of microstrip lines require an important area on the chip. The SWL technology can be a valuable candidate for minimizing the occupied area on the chip by the microstrip lines used for interconnection or phase shifting [START_REF] Tang; Franc | Performance Improvement Versus CPW and Loss Distribution Analysis of Slow-Wave CPW in 65 nm HR-SOI CMOS Technology[END_REF], [START_REF] Jin | A Slow-Wave Microstrip Line With a High-Q and a High Dielectric Constant for Millimeter-Wave CMOS Application[END_REF][START_REF] Dinc | An X-Band Slow-Wave T/R Switch in 0.25µm SiGe BiCMOS[END_REF].

The most common transmission lines implemented using SWL technology are the Shielded Coplanar Waveguide (S-CPW) [START_REF] Tang; Franc | Performance Improvement Versus CPW and Loss Distribution Analysis of Slow-Wave CPW in 65 nm HR-SOI CMOS Technology[END_REF]. Nevertheless Shielded MicroStrip Lines (S-MSL) can be designed. This particular type of SWL allows a very compact structure because a large gap is not required. This paper addresses the design the simulation and the experimental results of S-MSL implemented in a 0.25µm BiCMOS process from IHP.

Results:

The stack layers of the used BiCMOS process are represented in Fig. 1. The back end of this technology offers five metal layers: three thin aluminum layers and two thick layers fit for transmission lines and inductors design.

The position of the metal strips forming the ground plane varies for the two structures presented in this paper. The signal line is always located on the superior metal layer (TM2) while the periodic metal strips are located on the Top Metal 1 layer -TM1 (design 1) or, on the layer TM1 and M3 (design 2). The two thick metal layers TM2 and TM1 are separated by a 2 µm layer of SiO2. The capacitance per unit of length (C) is maximized by positioning the metal strips on the layers closest to the signal. However, even in this configuration, the value of the capacitance per unit of length is relatively small conducting to a poor quality factor. The metal strips in the second design are connected by vias (Fig. 2). This disposition of these strips prevents more effectively the E-field from penetrating inside the substrate and thus improves the quality factor Q.

The metal strips are disposed orthogonally to the direction of propagation so the return currents cannot flow directly under the signal contrary to classical micro-strips lines. By making these currents flow far away from the signal line, it is then possible to artificially improve the inductance per unit of length (L). A diminution of the phase velocity for a transmission line can be achieved by controlling C and L.

vp =c/ √ µ 𝑟𝑟 𝜀𝜀 𝑟𝑟 =1/√𝐿𝐿𝐿𝐿 (1)
where L is series inductance per unit length and C denotes shunt capacitance per unit length.

The structures (S-MSL and MSL) implemented in 0.25µm SiGe BiCMOS process were electromagnetically simulated using HFSS software. Manufactured samples were also measured by using a VNA (Anritsu ME7808A VNA). The S-parameters, attenuation (α) and phase (β) constants were measured/extracted after de-embedding. A special designed Thru, Reflect, Line (TRL) calibration kit was manufactured on the die for de-embed the impact of the microstrip access lines (length: 150 µm) on the S-MSL.

Fig. 3 Layout photograph of the manufactured die including a reference microstrip line, S-MSL and the TRL calibration kit

S-parameters measurements were performed with a VNA using a standard SOLT calibration procedure following by the TRL calibration. A layout view of the structures is shown on Fig. 3. The lines are 500 µm long (a minimum length of 300µm is required to avoid measurement errors due to the coupling between the probes). The reference MSL is also used as 'line element' in the TRL kit (Fig. 3 bottom left corner). The S-MSLs are preceded by a line length of Thru/2 that will be removed by de-embedding.

Fig. 4 shows the different measurements and the comparison between S-MSL and MS line. The width of the MSL signal line is w=15 µm conducting to a characteristic impedance of 50Ω for the reference MSL. The S-MSLs presented here were obtained after the optimization of three geometrical parameters: SL (the width of the metal strips shown Fig. 2), SS (the strip spacing Fig. 2) and the Gap (the distance between the signal line and the ground plane cf Fig. 1). The strip spacing was chosen to minimize the eddy currents loss (SS = 6µm, SL = 2µm). The experimental results of the phase constant β show that the structures are working for Ku-band. However the attenuation of the first S-MSL is 0.42 dB/mm (0.5 dB/mm for design 2) at 12.5 GHz versus 0.24 dB/mm for the MSL. The reference MSL have a quality factor inferior to S-MSL structures. The differences between MSL and S-MSL appear also on Fig. 4 (a).

The simulations results (not detailed here) demonstrated that closer to the signal is the ground plane more α and β decreases. The gap must be sized as function of the system specifications (losses and capacity). According to this work, a reduction about of 58% can be expected for the length of the transmission line by using a M-SWL instead of a classical MSL for a targeted phase shift. A 1.78mm long conventional MSL is necessary to achieve a 45° phase shift (insertion loss of 0.3-0.4dB). For the same phase shift a length of only 740 µm is required with the S-MSL (insertion loss of 0.25-0.34dB). The gap of the fabricated S-MSL is 50 µm and is relatively small compared to other SWL at 12.5 GHz.

Conclusion:

This paper presented two designs of S-MSL implemented in 0.25µm SiGe BiCMOS technology. The experimental results show an improvement of the phase constant by a factor 2.4 (β is improved by 2.4 at 12.5GHz) compared to the reference MSL while the attenuation constant is not strongly impacted. Thus the proposed S-MSLs are a very interesting candidate for phase shifting cells.

Fig. 2

 2 Fig. 2 Cross-section view of the S-MSL structure

Fig. 4

 4 Fig. 4 Comparison between MSL and S-MSL performances (measurements) a Attenuation constant (dB/mm) b Phase constant (rad/mm)

M

  . Gastaldi, D.Dragomirescu, A.Takacs (CNRS, LAAS, 7 avenue du colonel Roche, F-31400 Toulouse, France) E-mail:mgastald@laas.fr V.Armengaud (CNES RF/HT, 18 Avenue Edouard Belin, 31400 Toulouse, France) S.Rochette ( Thales Alenia Space, 26 av. JF. Champollion, BP1187, F-31037 Toulouse Cedex 1, France)