Haisheng Lin

Chengyong Wang
email: cywang@gdut.edu.cn

Zhiwei Lai

Tongchun Kuang

Mohamed Abdou Djouadi

Mohammed Abdou Djouadi

Microstructure and mechanical properties of HfBx coatings deposited on cemented carbide substrates by HiPIMS and DCMS

Keywords: HfB2 coating, magnetron sputtering, Microstructure, Hardness, Adhesion

Introduction

Hafnium diboride (HfB2) is a well-known ultra-high-temperature ceramic (UHTC) with a hexagonal structure. HfB2 has excellent physical properties, including a high melting point (3380 °C), high bulk hardness (28 GPa), high thermal conductivity (>100 W . (m . K) -1 ), low thermal expansion (6.3x10 -6 K -1 ) and good corrosion resistance [START_REF] Zhang | Thermal and Electrical Transport Properties of Spark Plasma-Sintered HfB2 and ZrB2 Ceramics[J][END_REF][START_REF] Gasch | Thermal conductivity characterization of hafnium diboride-based ultra-high-temperature ceramics[END_REF][START_REF] Fahrenholtz | Refractory Diborides of Zirconium and Hafnium[J][END_REF]. The lattice mismatch between HfB2 and GaN [START_REF] Yamada-Takamura | Surface and interface studies of GaN epitaxy on Si(111) via ZrB2 buffer layers[J][END_REF] is only 1.4%, suggesting that HfB2 is possibly an ideal buffer layer for the growth of semiconductors. Therefore, HfB2 coating is considered to have a high potential to be used in aerospace [START_REF] Upadhya | Advanced materials for ultrahigh temperature structural applications above 2000[END_REF][START_REF] Squire | Material property requirements for analysis and design of UHTC components in hypersonic applications[END_REF], thermophotovoltaic devices [START_REF] Arpin | Three-dimensional self-assembled photonic crystals with high temperature stability for thermal emission modification[END_REF], complementary metal oxide semiconductors (CMOS) [START_REF] Zagozdzon-Wosik | Applications of metallic borides for gate electrodes in CMOS integrated circuits[END_REF], and cutting tools. Hence, the study on the growth of high-quality HfB2 coating has excellent scientific and technological importance.

Obtained by chemical vapor deposition (CVD) from single-source precursor Hf(BH4)4 [START_REF] Yang | Crystalline texture in hafnium diboride thin films grown by chemical vapor deposition[END_REF][START_REF] Yang | CVD growth kinetics of HfB2 thin films from the single-source precursor Hf(BH4)(4)[END_REF][START_REF] Chatterjee | Nanoscratch and nanofriction behavior of hafnium diboride thin films[END_REF], the texture of HfB2 film is greatly dependent on the substrates like Si(111), Si(001), c-Si, SiO2, and deposition temperatures which are usually in the range of 400-900 °C [START_REF] Yang | Crystalline texture in hafnium diboride thin films grown by chemical vapor deposition[END_REF] and pressure [START_REF] Yang | CVD growth kinetics of HfB2 thin films from the single-source precursor Hf(BH4)(4)[END_REF]. Epitaxial HfB2 (0001) was grown on Si (001) substrate at a high temperature of 1000 °C by etching through a SiO2 layer [START_REF] Yang | Epitaxial growth of HfB2 (0001) on Si (001) by etching through a SiO2 layer[J][END_REF].

For the HfBx (x value is higher or lower than 2) coating deposited by radio frequencies magnetron sputtering (RFMS) [START_REF] Goncharov | Structure, composition, and physicomechanical characteristics of HfB2 and Hf-B-N films[END_REF][START_REF] Eldridge | Various properties of sputter-deposited HfB2 films[J][END_REF][START_REF] Yu | The influence of modulation periods on the evolution of microstructure and mechanical properties of nanoscale HfN/HfB2 multilayers[END_REF], its crystallization could be adjusted from the amorphous, cluster, poor pronounced columnar to columnar, with the variation of bias voltage and temperature simultaneously [START_REF] Goncharov | Structure, composition, and physicomechanical characteristics of HfB2 and Hf-B-N films[END_REF]. The corresponding hardness can be increased from 13 GPa for the amorphous film to 44 GPa for the (001) preferred oriented film [START_REF] Goncharov | Structure, composition, and physicomechanical characteristics of HfB2 and Hf-B-N films[END_REF]. A recent publication reported that the HfB2 coating deposited by DCMS on the Al2O3 substrate presents a hardness of 47.7 ± 2.7 GPa [START_REF] Glechner | Influence of the non-metal species on the oxidation kinetics of Hf, HfN, HfC, and HfB2 coatings[END_REF].

Moreover, HfB2 coating was deposited on stainless steel by DCMS with Al2O3 coating as a top layer [START_REF] Qiu | Structure, optical simulation and thermal stability of the HfB2-based high-temperature solar selective absorbing coatings[END_REF]. This HfB2/Al2O3 solar absorber coating was thermally stable in vacuum at 600 °C for 2 h, indicating that HfB2 could be a potential material for high-temperature solar thermal applications. Bakhit B et al. [START_REF] Bakhit | Dense Ti0.67Hf0.33B1.7 thin films grown by hybrid HfB2-HiPIMS/TiB2-DCMS co-sputtering without external heating[END_REF] reported the synthesis of dense Ti0.67Hf0.33B1.7 thin film with a hardness of ~41.0 GPa grown without external heating by hybrid HiPIMS and DCMS co-sputtering (HfB2-HiPIMS/TiB2-DCMS) on Al2O3 substrates. To the best of our knowledge, there are no reports on the deposition of HfBx coatings, including stoichiometric and off-stoichiometric compositions, on the cemented carbide substrate (WC/Co) as coated cutting tools. In addition, microstructure and mechanical properties of HfBx coatings deposited by magnetron sputtering technology, especially HiPIMS, also require further investigation.

In this paper, we adopted direct current magnetron sputtering (DCMS) and high-power impulse magnetron sputtering (HiPIMS) to deposit HfBx coating on cemented carbide substrates. The influence of different deposition parameters on the microstructure and mechanical properties of HfBx coating were studied. Furthermore, the adhesion mechanism between HfBx coating and cemented carbide was explored.

Experimental methods

Coating deposition

HfBx coatings were deposited on the mirror-polished (Ra < 0.04 μm) substrate of YL10.2 cemented carbide (manufacturer: ZCC•CT). Before the deposition, all the samples were ultrasonically cleaned and dried. The industrial-scale physical vapor deposition (PVD) equipment Domino Mini from Sulzer Metaplas, which is equipped with arc-enhanced glow discharge (AEGD), dcMS and HiPIMS technologies, was used for the deposition. The etching process with AEGD using a Ti target with 99.99 % purity was executed prior to the deposition of the coatings. The background pressure was evacuated to less than 4.0×10 -3 Pa before the deposition. HfB2 compound target (manufacturer: Plansee) was used to prepare coatings. The HfB2 target is 456 mm in length, 81 mm in width and 5 mm in thickness, and the purity of the target is 99.9 %. Magnets were mounted in a closed field configuration behind the target, and the effective coating area was 400 mm in length. Only one magnetron was used, and the samples were mounted in front of it. The minimum distance between sample and target is 87 mm. The working gas Ar had high purity of 99.999 %. The deposition conditions for different samples prepared using DCMS and HiPIMS are listed in Tables 1 and2, respectively. High power magnetron sputtering power supply was used in power control mode, with an average power of 4 kW. The monitoring screen showed that the discharge voltage and average current values were about 800 V and 5 A, respectively. The peak power was calculated to be 140 kW.

Characterization

A Bruker D8 Advance X-ray Diffractometer, with Cu K source and wavelength = 0.154 nm, was used to measure the phase structures and orientations of HfBx coatings. Hitachi SEM S-3400N-II acquired the surface morphology and cross-section structure of HfBx coatings. The acceleration energy of 15 kV was applied. To observe the cross-section structure of the coating, wire cutting was used to cut the coating samples from the cross-sectional direction. Bruker AFM Dimension FastScan Microscope was utilized to measure the 3D topology of the coating with the scanning area of 1 um×1 um. The microstructures of the HfBx coatings were observed by field-emission transmission electron microscopy (Thermo Fisher Scientific Talos F200S of FEI) with an EDS and an accelerating voltage of 200 kV. The sample used for transmission electron microscopy (TEM) detection was prepared by the focus ion beam. The surface chemical bonding states and the composition of the coatings were investigated using the X-ray photoelectron spectroscopy (XPS) with an Escalab 250Xi XPS Spectrometer (Thermo Fisher, USA). High-resolution spectra were acquired with a 0.1 eV step with a pass energy of 20 eV, corresponding to an overall Fermi edge resolution of 0.39 ± 0.02 eV. The binding energy of C1s (285.0 eV) was used to calibrate the energy scale of the spectra. A Shirley background subtraction method was beneficial in analyzing to differentiate the XPS peaks. Concerning the mechanical properties, Nano Indenter XP TM from Anton Paar was used to measure the hardness and Young's modulus of the coatings, with a depth of indention of 100 nm, dwelling time of 3 s and an applied load of 10 mN. Berkovich diamond indenter was used, with 25 indents per sample for nanoindentation measurement. The scratch tester RST from Anton Paar was used to evaluate the adhesion of HfBx coatings on cemented carbide substrates, with scratch loads ranging from 0 to 100 N, scratch length of 5 mm, and scratch speed of 10 mm/min. Rockwell indenter with a radius of 0.2 mm was used for the scratch measurement.

Results and discussion

Crystallinity and microstructure

The x-ray diffraction pattern of HfBx coatings deposited by DCMS is shown in Fig. 1. Heating temperature and bias voltage were changed and the corresponding XRD patterns were obtained. The peak around 24.5°is assigned to HfB2 with (001) orientation, and it has the highest intensity relative to all the other peaks. Peaks from HfB2(100), (101), (002) and (102) orientations are also found and marked in Fig. 1, which indicate the polycrystalline nature of the HfBx coatings. With the increase of bias voltage from -50 to -200 V at a heating temperature of 600 °C, the peak intensity associated with (001) orientation becomes higher. In addition, the peak intensity associated with (001) orientation corresponding to a low heating temperature (500 °C) and a high bias voltage (-150 V) is lower than the peak intensity corresponding to a high temperature (600 °C) but a low bias voltage (-100 or 50 V). It implies that heating temperature and bias voltage are critical factors for the crystallinity of HfBx coatings with (001)-preferred orientation using DCMS. It is also consistent with the growth principle of the coatings proposed in the literature [START_REF] Thornton | The microstructure of sputter-deposited coatings[END_REF]. Moreover, the HiPIMS-deposited HfBx coating crystals are highly oriented in the HfB2(001) orientation along the c axis. Likewise, the nanocolumnar structure along the direction of growth of the coating is also found and shown in Fig. 3(b).

HfBx coating with HfB2(001)-preferred oriented can be obtained regardless of the sputtering methods like DCMS, HiPIMS and RFMS [START_REF] Buranich | Mechanical and tribological characterization of nanostructured HfB2 films deposited from compound target[END_REF]. The texture evolution of HfBx coatings can be explained by adatom surface diffusion [START_REF] Greene | Development of preferred orientation in polycrystalline TiN layers grown by ultrahigh vacuum reactive magnetron sputtering[END_REF], which occurs among planes and among grains during the initial or growth stages of the surface of films [START_REF] Kajikawa | Comprehensive perspective on the mechanism of preferred orientation in reactive-sputter-deposited nitrides[J][END_REF]. It greatly influenced the texture evolution of other transition metal diborides like TiB2 [START_REF] Zhang | Influence of negative bias voltage and deposition temperature on microstructure and properties of superhard TiB2 coatings deposited by high power impulse magnetron sputtering[END_REF] and CrB2 [START_REF] Zhang | Temperature induced superhard CrB2 coatings with preferred (001) orientation deposited by DC magnetron sputtering technique[END_REF]. According to the theory reported in [START_REF] Kajikawa | Comprehensive perspective on the mechanism of preferred orientation in reactive-sputter-deposited nitrides[J][END_REF], in the case of low surface adatom mobility, grains on the plane that has lower surface diffusivity parallel to the substrate survive and become the preferred orientation. However, at higher surface adatom mobility and longer diffusion distances, diffusion among grains occurs and grains with the lowest-surface-energy plane become the preferred orientation.

Concerning hexagonal HfB2 (h-HfB2), Zhang et al. calculated the surface energy, with thermodynamic sense of several HfB2 surfaces, including (0001), (10 0), (10 1) and ( 103). They found that Hf-terminated (0001) surface and B-terminated (0001) surface have the lowest surface formation energy at the Hf-rich area and B-rich end area, respectively [START_REF] Zhang | First-principles investigation of the thermodynamic stability of MB2 materials surfaces (M=Ti/Zr/Hf)[END_REF]. At higher deposition temperatures, atoms gain more thermal energy and consequently have higher adatom mobility. The growing coating surface gains additional energy from high ion bombardment by increasing bias voltage.

Consequently, the coating favors the growth of the lowest-surface-energy orientation (001). Due to this reason, a highly HfB2(001)-oriented coating is obtained after using high deposition temperature and bias voltage in DCMS. Compared to DCMS, HiPIMS-deposited HfBx coating exhibits a higher intensity of the peak associated with (001) orientation, but has fewer peaks assigned to other orientations. It is because of the continuous ion bombardment on the coating surface from the high ionization rate of HiPIMS [START_REF] Ehiasarian | Time-Resolved Ionisation Studies of the High Power Impulse Magnetron Discharge in Mixed Argon and Nitrogen Atmosphere[END_REF]. It leads to high surface adatom mobility [START_REF] Mayrhofer | A comparative study on reactive and non-reactive unbalanced magnetron sputter deposition of TiN coatings[END_REF] and highly HfB2(001)-oriented HfBx coating.

Surface and cross-sectional morphology

SEM and AFM were used to observe the cross-sectional view, surface morphology and 3D-topology of different HfBx coatings deposited by DCMS and HiPIMS, as shown in Fig. 4 and5.

HfBx coatings deposited by DCMS, with different deposition conditions, have a prominent columnar cross-sectional structure, which is of coarse and have low density.

On the other hand, all HfBx coatings deposited by HiPIMS show a highly compact and smooth structure without macro features.

Based on the surface observed by SEM and AFM, HfBx coatings deposited by DCMS have small spherical islands. Their arrangement on the surface is random with non-uniform size. Some small "islands" get together to form an "island clusters", and a clear boundary of among "island clusters" can be observed. However, HfBx coatings deposited by HiPIMS have an extremely smooth surface, and no particles can be observed in the magnified view, as shown in Fig. 5.

HfBx coatings with significantly different morphologies were obtained by DCMS and HiPIMS. Due to the high ionization and high peak power density in HiPIMS, surface adatom mobility is greatly improved compared to conventional DCMS technology, resulting in a dense and uniform structure. In addition, high ion bombardment in HiPIMS contributes to grain refining and prevents from the columnar structure formation .

Mechanical properties

Hardness and Young's modulus of HfBx coatings deposited by DCMS are shown in Fig. 6. For the coatings deposited at low heating temperature (500 °C), relatively low hardness and Young's modulus are measured with values of 27.3 ± 4.2 and 470.5 ± 6.6 GPa, respectively. It is because of the low crystallinity of the coatings. For the HfBx coatings deposited at higher heating temperature (600 °C), with the increase of bias voltage from -50 to -200 V, the hardness increases slightly from about 33.4 ± 2.2 to 36.3 ± 2.5 GPa. For the same increase of bias voltage, Young's modulus increases from 470.5 ± 6.6 to 567.4 ± 8.1 GPa. It is due to higher ion bombardment on the coating surface during the deposition and formation of a structure with higher crystallinity, as shown in the XRD results in Fig. 1.

Considering the deposition by HiPIMS, all the HfBx coatings have superhardness (≥ 40 GPa), illustrated in Fig. 7. By increasing the bias voltage from -100 to -200 V (at working pressure 0.6 Pa), the hardness and Young's modulus can be increased from 42.6 ± 3.8 to 45.1 ± 2.5 GPa, and 571.8 ± 8.9 to 583.0 ± 7.0 GPa, respectively. Additionally, on decreasing the working pressure from 0.65 to 0.5 Pa, the hardness and Young's modulus of HfBx coating further increase from 45.1 ± 2.5 and 583.0 ± 7.0 GPa to 49.3 ± 3.6 and 667.0 ± 9.2 GPa, respectively. The different values of hardness for HfBx coatings under different conditions are consistent with the crystallinity, as per the XRD results shown in Fig. 2.

From the first principle calculations, h-HfB2 is a material of hard nature with a calculated Vickers hardness of 40.7 GPa [START_REF] Pan | Phase stability and mechanical properties of hafnium borides: A first-principles study[END_REF] and the weakest ideal strength of 33.3 GPa [START_REF] Zhang | The Ideal Strength of Transition Metal Diborides TMB2 (TM = Ti, Zr, Hf): Plastic Anisotropy and the Role of Prismatic Slip[J][END_REF]. Note that the theoretically calculated bulk hardness is based on the ideal crystal structure of h-HfB2. The mechanical properties of the experimentally deposited coatings are greatly influenced by crystallinity and microstructure when deposited by different processes. For example, a previous study [START_REF] Fuger | Anisotropic super-hardness of hexagonal WB2±z thin films[END_REF] reported that nanoindentation revealed 15 GPa higher hardness of (0001)-oriented h-WB2 coating as compared to the (1011)-orientated one.

According to the hardness measuring results in this study, HfBx coating can be hard or even superhard. As for the coatings deposited by HiPIMS, the ion flux is much larger than other sputtering methods, such as DCMS and middle-frequency pulsed MS [START_REF] Ehiasarian | Time-Resolved Ionisation Studies of the High Power Impulse Magnetron Discharge in Mixed Argon and Nitrogen Atmosphere[END_REF]. High atomic mobility can be induced by bombarding the growing coating with energetic heavy Hf+ ions generated by HiPIMS sputtering of the HfB2 target [START_REF] Bakhit | Dense Ti0.67Hf0.33B1.7 thin films grown by hybrid HfB2-HiPIMS/TiB2-DCMS co-sputtering without external heating[END_REF]. According to the present study, the deposition rate of the HiPIMS-deposited HfBx coating (about 5 nm/min) is much lower than the DCMS-deposited one (about 11~14 nm/min). When HiPIMS is used, the coating structure is dense and has no micro features, and the coatings would have higher crystallinity in (001) orientation as compared to DCMS. HfBx coatings deposited by DCMS have low crystallinity, random orientations and coarse columnar structure, due to the fast deposition rate and insufficient energy of particles upon reaching the coating surface. Therefore, the coatings prepared by HiPIMS exhibits higher hardness than those prepared by DCMS.

In order to explore the effect of the composition ratio in HfBx coating on the mechanical properties of the coating, high-resolution XPS B1s spectra of HfBx coatings was observed, as shown in Fig. 8. The peaks at about 192.1, 187.75 and 186.7 eV are attributed to B in the B2O3 [START_REF] Burke | Ignition mechanism of the titanium-boron pyrotechnic mixture[END_REF], HfB2 [START_REF] Chowdhury | Nanomechanical and nanotribological behaviors of hafnium boride thin films[END_REF] and metallic boron [START_REF] Shea | Springer handbook of condensed matter and materials data[END_REF][START_REF] Nedfors | Influence of pulse frequency and bias on microstructure and mechanical properties of TiB2 coatings deposited by high power impulse magnetron sputtering[END_REF][START_REF] Hendrickson | Core-electron binding energies for compounds of boron, carbon, and chromium[END_REF][START_REF] Schreifels | X-ray photoelectron spectroscopy of nickel boride catalysts: correlation of surface states with reaction products in the hydrogenation of acrylonitrile[END_REF] respectively. Note that air-exposed samples were used for the measurement and sputtering process was avoided. It was because of the risk of preferential removal or implantation of low atomic mass species (boron) by sputtering, making accurate phase identification difficult (impossible). Although most TMB2 coatings are over-stoichiometric [START_REF] Mayrhofer | Self-organized nanocolumnar structure in superhard TiB2 thin films[END_REF][START_REF] Audronis | A TEM study of the structure of magnetron sputtered chromium diboride coatings[END_REF][START_REF] Ge | Hard and wear resistant VB2 coatings deposited by pulsed DC magnetron sputtering[END_REF], some have been reported to be under-stoichiometric [START_REF] Hellgren | High-power impulse magnetron sputter deposition of TiBx thin films: Effects of pressure and growth temperature[END_REF][START_REF] Thörnberg | Microstructure and materials properties of understoichiometric TiBx thin films grown by HiPIMS[END_REF]. However, the presence of the B-B peak in XPS indicates that these HfBx coatings are indeed over-stoichiometric. Possible reasons, which need to be verified, may include the longer mean free path of the B atom than the Hf atom, or the much higher mass of Hf compared to Ar resulting in slight deflection of Hf. In the study of the compositional evolution of Ti-B thin films deposited by sputtering a compound target, Neidhardt J et al. found that the coatings composition Ti/B ratio mainly depends on gas pressure, target-substrate distance and angle [START_REF] Neidhardt | Experiment and simulation of the compositional evolution of Ti-B thin films deposited by sputtering of a compound target[END_REF]. HfBx coatings, with constituents with a large difference in atomic mass or radius, may also have similar laws of influence [START_REF] Neidhardt | Experiment and simulation of the compositional evolution of Ti-B thin films deposited by sputtering of a compound target[END_REF], and further research studies are needed to verify.

The B-B/B-Hf ratios is calculated from the area of B1s spectra, as shown in Fig. 9.

The calculated results show that the B-B/B-Hf ratio of HiPIMS-deposited coating is higher than that deposited by DCMS. Considering the higher hardness of HiPIMS-deposited HfBx coatings as compared to that deposited by DCMS, a higher amount of boron may be associated with higher hardness.

The mechanism of the enrichment or deficiency of the boron phase on the hardening of the HfBx coating needs further deep investigation. The contribution of B-B bonding from B1s spectra was assigned to the thin B-tissue phase encapsulating the TiB2 nanocrystals observed by TEM [START_REF] Nedfors | Influence of pulse frequency and bias on microstructure and mechanical properties of TiB2 coatings deposited by high power impulse magnetron sputtering[END_REF]. In the present study of HfBx coatings, a similar structure to the TiB2 coating with B-enrichment [START_REF] Mayrhofer | Self-organized nanocolumnar structure in superhard TiB2 thin films[END_REF] was observed by HRTEM (Fig. 3). Excess B phase tissue along the TMB2 column was also found from the study of other TMB2 coatings, i.e., TiB2 [START_REF] Mayrhofer | Self-organized nanocolumnar structure in superhard TiB2 thin films[END_REF], CrB2 [START_REF] Audronis | A TEM study of the structure of magnetron sputtered chromium diboride coatings[END_REF] and VB2 [START_REF] Ge | Hard and wear resistant VB2 coatings deposited by pulsed DC magnetron sputtering[END_REF]. Mayrhofer et al. [START_REF] Mayrhofer | Self-organized nanocolumnar structure in superhard TiB2 thin films[END_REF] and N. Kalfagiannis et al. [START_REF] Kalfagiannis | Excess of boron in TiB2 superhard thin films: a combined experimental and ab initio study[END_REF] associated the hardening effect of the coating with the formation of these B phases. It is reported that higher crystallinity in (001) orientation for RFMS-deposited HfBx coating has a higher amount of B/Hf ratio and higher hardness [START_REF] Goncharov | Structure, composition, and physicomechanical characteristics of HfB2 and Hf-B-N films[END_REF]. HfBx coatings might also be under-stoichiometric. In the case of under-stoichiometric TMBx [START_REF] Hellgren | High-power impulse magnetron sputter deposition of TiBx thin films: Effects of pressure and growth temperature[END_REF][START_REF] Thörnberg | Microstructure and materials properties of understoichiometric TiBx thin films grown by HiPIMS[END_REF], recent publication reports that deficiency of B in under-stoichiometric HiPIMS-deposited TiB1.43 film is found to be accommodated by Ti-rich planar defects, and the film exhibit a superior hardness of 43.9 ± 0.9 GPa [START_REF] Thörnberg | Microstructure and materials properties of understoichiometric TiBx thin films grown by HiPIMS[END_REF].

The scratch on the surface of the DCMS-deposited coating is shown in Fig. 10a.

Minimal chipping of the coating suddenly occurs at the side of the scratch track at the critical load of 78.5 N. No apparent delamination or chipping of the coating from the surface can be observed with scratching from 0 to 100 N, indicating excellent adhesion of the coating to the substrate.

As shown in Fig. 10b, there is no cohesive failure (i.e., failure of the coating before delamination) on the surface of HiPIMS-deposited coating under the scratch loads from 0 to 100 N, suggesting excellent adhesion strength of the coating. Tiny localized hemispherical chipping of the coating begins to appear at the side of the track at the critical load of 55 N. However, the coating removal inside the scratch track exacerbates smoothly with the load. It is in contrast to the DCMS-deposited coating, where chipping of the coating occurs suddenly. It is due to higher damage resistance and higher hardness of the coating deposited by HiPIMS.

In order to profoundly investigate the adhesion behavior between the coating and cemented carbide (WC/Co) substrate from a microscopic view, HRTEM was conducted to observe the interface, as shown in Fig. 11 and 12. The HfBx coating was deposited directly on the substrate without a buffer layer. Figure 11 shows the interface of HfBx coating and WC grains. The coating and the WC grain are perfectly connected. EDS was also conducted on the interface, and the inset in Fig. 11 is the line-scan EDS signal scanned from the coating toward the substrate. It is found that both of the heavy elements Hf and W have gradient diffusion towards the other side.

The WC grains are followed by a thin amorphous layer of 2-3 nm thickness before the formation of HfB2 crystals. It can also be observed that the crystallinity of HfB2 is low at the start of the nucleation stage.

Due to the etching on the surface of the cemented carbide substrate before deposition and different etching rates of WC grains and Co phases, a larger amount of Co is removed than WC grains. Therefore, a "nano-rough" surface with Co phases at lower positions is obtained. The growth of coating in such an area with surfaces of different heights contributes to the mechanical "locking" of the coating, as shown in Fig. 12, and improves the adhesion of the coating on the substrate.

Consequently, element diffusion and mechanical "locking" on the interface lead to excellent adhesion of HfBx coating on cemented carbide substrate.

Conclusion

In this study, HfBx coatings were deposited on cemented carbide substrate using DCMS and HiPIMS. The effects of the deposition process on the crystal structure, coating morphology and mechanical properties were analyzed, leading to the following results.

Higher heating temperature, higher bias voltage and lower working gas pressure, contribute to more thermal energy and kinetic energy on the surface, leading to higher adatom mobility and HfBx coating with stronger HfB2(001)-preferred orientation.

Moreover, high ion bombardment of HiPIMS results in the refining of grains, dense cross-sectional microstructure without obvious micro-feature, and smooth surface morphology, whereas the dcMS deposited HfBx coating has a columnar structure and surface morphology with small non-uniform spherical-like islands.

For both sputtering methods, the hardness of HfBx coatings increases with crystallinity. HiPIMS-deposited HfBx coatings possess a maximum superhardness up to a value of 49.3 ± 3.6 GPa. It is because of the coating's hard nature and high crystallinity. The excellent adhesion strength of the coatings on the cemented carbide substrates, without cohesive failure, can be obtained for both DCMS and HiPIMS methods. Further investigation reveals that such an excellent adhesion strength originated from the thin amorphous layer with diffusing elements in the interface and the mechanical locking due to the step structure among coating, WC, and Co grains.

Figure 2

 2 Figure 2 shows the XRD pattern of HfBx coatings deposited by HiPIMS with different bias voltages and argon gas pressure. It can be found that the highest peak of the HfB2 in the XRD pattern corresponds to the (001) orientation. Unlike the XRD pattern of HfBx coating deposited by DCMS, fewer peaks can be observed from other orientations in the XRD pattern of HfBx coatings deposited by HiPIMS. The HfBx coatings with HfB2(001) orientation and a bias voltage of -100 V possess a peak at 25.0°. The same peak shifts to 25.17°with a bias voltage of -200 V. The shift of the peak indicates higher compressive stress of the HfBx coating. It is due to stronger ion bombardment on the coating surface because of higher bias voltage. No noticeable difference in the intensity of (001) orientation peaks can be observed for different bias voltages regarding HfBx coatings deposited by HiPIMS. Moreover, when the working pressures of argon gas changed from 0.65 to 0.5 Pa, a change in intensity and slight shift of the angle to 25.22°of (001) orientation peaks can be found. Lower working pressure contributes to a smaller mean free path and reduced energy loss by the collision of particles in plasma. Particles reaching the surface of the coating with higher kinetic energies induce better growth of HfB2 grains. The cross-sectional high-resolution transmission electron microscope (HRTEM) shows that the DCMS-deposited HfBx coatings possess multiple orientations like HfB2(001), (100), (101) and others, indicating the polycrystalline nature of the coating.

Fig. 1 XFig. 2 XFig. 3 Fig. 4

 1234 Fig. 1 X-ray diffraction pattern of HfBx deposited by DCMS (Deposition parameters: 4 kW, 160 sccm Ar, 0.78 Pa.)

Fig. 5

 5 Fig. 5 Cross-section, surface morphology and 3D topology of HfBx coating deposited by HiPIMS with different deposition conditions. (a-c) HiP-1 sample; (d-f) HiP-2 sample; (g-i) HiP-3 sample.

Fig. 6 Fig. 7 Fig. 8 Fig. 9 Fig. 11 Fig. 12

 67891112 Fig. 6 Hardness and Young's modulus of HfBx coating deposited by DCMS (Deposition parameters: 4 kW, 160 sccm Ar, 0.78 Pa.)

Table 1

 1 Deposition conditions of various DCMS-deposited HfBx coatings

	Sample	DC-1	DC-2	DC-3	DC-4
	Power (kW)		4
	Ar (sccm)		160
	Pressure (Pa)		0.78
	Bias voltage (V)	-150	-50	-100	-200
	Heating temperature (°C)	500	600	600	600
	Deposition time (min)		150

Table 2

 2 Deposition conditions of various HiPIMS-deposited HfBx coatings

	Sample	HiP-1	HiP-2	HiP-3
	Average Power (kW)		4
	Ar (sccm)		140
	Pressure (Pa)	0.65	0.65	0.5
	Bias voltage (V)	-100	-200	-200
	Heating temperature (°C)		600
	On/off (us)		40/1000
	Deposition time (min)		300

Acknowledgments

This work was supported by the National Natural Science Foundation of China Youth Project (Grant no. 52005111).

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.


The slurry instability is a serious and well-known issue affecting the industrial production of lithium ion batteries and is caused by residual lithium compounds present at the surface of the cathode material (Li(Ni x Mn y Co 1-x-y )O 2 or NMC). Although the problem is recognized by the battery industry, the scarce communication between academia and industry and the limited access to industrial information lead to an important disparity in the knowledge of this issue between the two realities. 1 The often observed "slurry gelation" is most likely due to the PVdF instability and leads to catastrophic issues in the mass production line. An optimal viscosity of the slurry is actually essential to obtain high quality slurries and thence high performance batteries. Residual lithium species are usually considered as limited to LiOH and Li 2 CO 3 . 1 Consequently, a surface reaction between the more basic LiOH and PVdF is frequently suggested in the literature as a possible cause of the widely observed slurry gelation. [1][2][3][4][5][6] Such a phenomenon was indeed found to occur more likely for high Ni compositions -more affected by the formation of lithiated surface species-than during the preparation of electrodes based on NMCs with lower Ni content. According to this explanation, the PVdF would undergo dehydrofluorination following the simplified reaction (CH 2 -CF 2 ) n + LiOH  (CH=CF) n + LiF + H 2 O proposed by G.J. Ross et al. 7 However, as highlighted by Seong et al., this topic has not been critically considered in the literature yet and a proper method to confirm the suggested reaction and to accurately characterise the residual lithium species has never been used to this end. In this communication, we show how a detailed description of the relevant compounds, and their quantification, can be obtained by using a combination of 7 Li MAS-NMR and XPS. The analysis of industrial powder and slurry samples was conducted to suggest a new reaction for the PVdF binder degradation, involving Li 2 O as the reagent. Three industrial NMC811 powders (UMICORE) with different amounts of surface lithiated species and the corresponding dried slurries (mixtures of active material (AM), conductive carbon, PVdF and NMP) were analysed (Table 1). The powders were obtained placing the pristine industrial powder under three controlled atmospheres, allowing to attain different surfaces and reactivities, as highlighted in a previous study. 8 Sample 1 and Sample 2 powders were obtained after aging in air under relative humidity RH=30% for 15 minutes and 2 days, respectively. On the other hand, the pristine powder was left in synthetic air atmosphere with RH ≈ 70% and without any carbon dioxide for 2 days to prepare Sample 3. Dried slurries were then made at Umicore's production plant using these three powders as AM, in typical fabrication conditions. The analysis of dried slurries instead of electrodes coated on the current collector was performed to avoid energetic processes inducing further degradation of the polymer. 7 Li MAS-NMR was used to estimate the amount of surface diamagnetic lithium, as described in our previous article. 8 Results are reported in Table 1. The extraction of lithium from the AM to form the considered surface species increases in all the samples during the slurry preparation step, demonstrating that additional changes affect the NMC811 surface upon this procedure. 1. Samples description. Surface diamagnetic lithium amounts evaluated by 7 Li MAS-NMR given in µmol Li /g AM with an error of ±10%

Amount (µmol

The analysis of XPS spectra evidences the presence of LiF in all the analysed slurries. This compound, obviously absent in the powders, is characterised by a peak around 684.5 eV; this component is visible in the typical F 1s spectrum, obtained for these samples, reported in Figure 1a. 19 F MAS-NMR was also performed to confirm the presence of LiF in the dried slurries independently of the X-Ray exposure, which was proven to be energetic enough to break C-F bonds in too harsh conditions. 9 Figure 1b represents a typical spectrum obtained for slurries, showing the peak around -205 ppm, corresponding to the LiF contribution, indicated by straight orange arrows (see the inset of Figure 1b). Its relative spinning sidebands (dashed orange arrows in Figure 1b) are also marked. Since the slurry is not exposed to any energetic source during MAS-NMR analysis, the hypothesis of a PVdF degradation coming from X-rays exposure can definitely be excluded and new hypotheses to explain the appearance of an LiF signal are needed. The identification of LiF allowed us fully deconvoluting the isotropic resonance of the 7 Li MAS-NMR spectra by including a contribution at -1 ppm (Fig. 2, complete spectra and deconvolutions in Fig. S1) in addition to the four contributions (Li 2 CO 3 , LiOH, Li 2 SO 4 , Li 2 O) already confirmed for the pristine powder as detailed in our previous work. 8 The method combining NMR and XPS data is explained in Reference 8 . It was thence possible to estimate the amount of each species (in µmol/g AM ) found in the six (powders and slurries) samples (also collected in Table S1 in SI together with the parameters used for the fit). Before proceeding to the interpretation of these results, some preliminary observations are needed: i-Slurries are prepared in an industrial dry room. The relative humidity of this environment (typically controlled to be between 0.7 and 1.5%) is sensibly lower than that of the climatic chamber samples were aged in. Reactions with H 2 O (if there is any occurring) will be therefore ignored when analysing the effect of the electrode preparation step. Conversely, the presence of elevated amounts of CO 2 in the dry room suggests that reactions with the carbon dioxide should be considered. ii-Li 2 SO 4 coming from the synthesis (residue from the sulphate precursors) is considered to be inhomogeneous in the analysed samples. Any increase or decrease in its quantity is thus not taken into account. iii-In all cases, the following order is found for the 7 Li NMR FWHMs: Li 2 O> LiOH> Li 2 CO 3 , suggesting that, overall, Li 2 is closest to the active material while Li 2 CO 3 is mostly located in the outermost part of the surface contamination layer. 10 After the slurry preparation step, the LiF FWHM is always larger than those for LiOH and Li 2 CO 3 (with exception of Sample 1 in which FWHM LiF = FWHM LiOH ), suggesting a stronger interaction with the AM. Hence, the formation of LiF would occur at the interface between Li 2 O (very close to the AM) and LiOH (closer to the material than Li 2 CO 3 ). Taking into account these arguments and the quantities reported in Table S1, we suggest that the degradation of the PVdF binder to form the observed LiF is induced by Li 2 O, and not LiOH, as previously proposed in the literature, according to the reaction (CH 2 -CF 2 ) n + Li 2 O  (CH=CF) n + LiOH + LiF (Reaction 1). Moreover, considering that Li 2 O is detected in a considerably higher amount than LiOH by 7 Li MAS-NMR, its involvement in the degradation of the polymer binder appears to be most reasonable. As a matter of fact, the LiOH quantity increases in both Sample 2 and 3 slurries, in apparent contradiction with its consumption upon reaction with the PVdF to give LiF. Moreover, Li 2 O (or lithia) is an extremely basic species, which could thence react with the PVdF more easily than LiOH, due to its higher basicity. 11 In addition, it is particularly convincing to compare the F 1s XPS spectra acquired by Im et al. on pristine electrodes based on pure lithia, and those presented by Grissa et al. for NMC532 PVdF based electrodes. 4,6 NMC electrodes appear significantly less affected by PVdF degradation than lithia ones, even after heavy calendering of the former. 4,6 A quantitative examination of the NMR data can help confirming the implication of lithia in the degradation process. We can now follow the evolution of the amount (in µmol/g AM ) of the species taking part in the suggested Reaction 1, as found by 7 Li MAS-NMR, and reported in Table S1. It should also be reminded that due to the constant presence of carbon dioxide, LiOH is consumed to form Li 2 CO 3 via the reaction: 2LiOH + CO 2  Li 2 CO 3 + H 2 O (Reaction 2). Let us take Sample 2 as an example. During the electrode preparation (slurry), 3 µmol/g AM of LiF are formed at the NMC811 surface. According to Reaction 1, this should be accompanied by the formation of the same quantity of LiOH. Similarly, 1 µmol/g AM of Li 2 CO 3 was formed during the process, as indicated in Table S1, meaning that 2 µmol of LiOH were consumed, according to Reaction 2. Consequently, the amount of LiOH should increase by 1 µmol/g AM (3 µmol/g AM formed -2 µmol/g AM consumed), which is in line with the quantities found by 7 Li MAS-NMR. By using the same reasoning, the LiOH evolution upon electrode preparation of Sample 3 can be perfectly explained as well, confirming our hypothesis.

Concerning Sample 1, the evolution of lithiated species amounts is not completely in line with previous observations for Sample and 3. Actually, following the above explained process, we found that 0.4 µmol of LiOH should have been consumed during the slurry preparation step, whilst 1 µmol are actually consumed according to results reported in Table S1. Since Sample 1 was aged only 15 minutes in the climatic chamber, the NMC811 surface had most likely not completely reacted with the humid atmosphere, making the assumption (i) above mentioned untrue for this sample. If the AM surface were not in equilibrium with water contained in the atmosphere, multiple reactions could still occur during the slurry preparation. Therefore, the identification of reactions taking place at the Sample 1 surface turns out to be more complicated than for Samples 2 and 3, emphasizing the complexity of reactions occurring at the surface and involving the surface of the active material, the surface species, and the surrounding media

Residual

  Li 2 O degrades PVdF during the preparation of NMC811 slurries for Li-ion batteries Angelica Laurita, *a Liang Zhu, b Pierre-Etienne Cabelguen, b Jérémie Auvergniot, b Dominique Guyomard, a Philippe Moreau a and Nicolas Dupré *a a Université de Nantes, CNRS, Institut des Matériaux Jean Rouxel (IMN), F -44000, Nantes, France b Umicore, 31 rue du marais, Brussels BE-1000, Belgium. * : Corresponding authors 7 Li MAS NMR quantification of lithiated species at the surface of aged NMC811 industrial powders and slurries shows that the electrode preparation process exacerbates Li extraction. A combination of 7 Li MAS NMR and XPS suggest a new reaction for the PVdF binder degradation, involving in fact Li 2 O as the reagent and the formation of LiF.

Figure 1 .

 1 Figure 1. Experimental proof of the occurrence of LiF (in orange) at NMC811 dried slurry surfaces by XPS (a) and 19 F MAS-NMR (b).

Figure 2 .

 2 Figure 2. Deconvolution of 7 Li MAS NMR spectra of NMC811 powder (a) and dried slurry (b)

  . The reaction path by which Li 2 O induces the PVdF degradation is hard to infer from our study, but it would be surely different from that relying on the high nucleophilic nature of OH -ions, mainly proposed in aqueous solutions. We believe that the liquid media, be it water or minute H 2 O quantities often present in NMP, are not essential to the occurring of the reaction. As a matter of fact, it was proven that positive electrodes calendering (i.e. in dry state) greatly increases the LiF production; similarly, such effect is also observed by Im et al. in lithia based electrodes. 4 In summary, we propose a new interpretation of LiF production in NMC811 slurries based on MAS-NMR and XPS results, induced by the presence of reactive Li 2 O at the surface of the positive electrode material. Neither the quantity of LiOH at the AM surface nor the minute quantities of water in NMP would be decisive parameters in the degradation of the PVdF binder. Lithia quantity should thus be considered more carefully in order to preserve the mechanical properties of these high-energy Li-ion battery electrodes.


J Trahan

J Profili

G Robert-Bigras

M Mitronika

M Richard-Plouet

L Stafford

Optical response of plasmonic silver nanoparticles after treatment by a warm microwave plasma jet

This work investigates the effect of plasma treatment on the morphology and composition of 15 x 15 mm 2 silver nanoparticle (70-80 nm) thin films. The silver nanoparticles are deposited onto thermal silica (SiO 2 /Si) substrates by spin-coating, then they are treated by an open-to-air microwave argon plasma jet characterized by a neutral gas temperature of 2200 ± 200 K.

Scanning electron microscopy analysis reveals that the number of isolated nanoparticles in the film sample decreases after exposure to multiple jet passes, and that polygonal structures with sharp corners and edges are produced. Similar structures with much rounder edges are obtained after conventional thermal annealing at temperatures up to 1300 K. Based on localized surface plasmon resonance (LSPR) analysis in the range of 350-800 nm, the main extinction band of silver nanoparticles experiences a redshift after treatment with the plasma jet or with thermal annealing. Moreover, both treatments induce surface oxidation of the nanoparticles, as evidenced by X-ray photoelectron spectroscopy. However, only the plasmaexposed samples exhibit a significant rise in the surface-enhanced Raman scattering (SERS) signal of oxidized silver at 960 cm -1 . 29 29 µm 2 mappings of hyperspectral Raman IMAging (RIMA) and multivariate curve resolution analysis by log-likelihood maximization demonstrate that the SERS signal is controlled by large-scale micrometer domains that exhibit sharp corners and edges.

Introduction

In light of their high thermal conductivity, excellent electronic conductivity [START_REF] Alshehri | Enhanced electrical conductivity of silver nanoparticles for high frequency electronic applications[END_REF][START_REF] Huang | Properties of polyacrylic acid-coated silver nanoparticle ink for inkjet printing conductive tracks on paper with high conductivity[END_REF][START_REF] Polavarapu | Preparation of conductive silver films at mild temperatures for printable organic electronics[END_REF], and well-established antimicrobial properties [START_REF] Dankovich | Bactericidal paper impregnated with silver nanoparticles for point-of-use water treatment[END_REF][START_REF] Fan | Chemical, electrochemical, gravimetric, and microscopic studies on antimicrobial silver films[END_REF][START_REF] Jamuna-Thevi | Quantification of silver ion release, in vitro cytotoxicity and antibacterial properties of nanostuctured Ag doped TiO 2 coatings on stainless steel deposited by RF magnetron sputtering[END_REF][START_REF] Kim | Synthesis and characterization of antibacterial Ag-SiO 2 nanocomposite[END_REF], silver-containing materials are used in a variety of applications. Silver-based thin films may be prepared using multiple processes, including vapor deposition and ion implantation [START_REF] El-Nour | Synthesis and applications of silver nanoparticles[END_REF][START_REF] Carles | Three dimensional design of silver nanoparticle assemblies embedded in dielectrics for Raman spectroscopy enhancement and dark-field imaging[END_REF], as well as nebulization, inkjet printing, and spin-coating of Ag nanoparticle (NP) dispersions [START_REF] Hong | Controlled two-dimensional distribution of nanoparticles by spin-coating method[END_REF][START_REF] Prevo | Controlled, rapid deposition of structured coatings from micro-and nanoparticle suspensions[END_REF]. When sintered by conventional thermal annealing, Ag NPs can exhibit very good electronic conductivity, which is relevant for a number of devices, including photovoltaic cells, light-emitting diodes, organic thin film transistors, and sensors [START_REF] Ahn | Omnidirectional printing of flexible, stretchable, and spanning silver microelectrodes[END_REF][START_REF] Shen | Preparation of solid silver nanoparticles for inkjet printed flexible electronics with high conductivity[END_REF]. For antimicrobial applications, the synthesis of Ag NPs with stable and defined shape, size, and chemical bonds is needed to achieve high chemical and biological activity [START_REF] Balout | Density functional theory study of the spontaneous formation of covalent bonds at the silver/silica interface in silver nanoparticles embedded in SiO 2 : Implications for Ag + release[END_REF][START_REF] Morones | The bactericidal effect of silver nanoparticles[END_REF][START_REF] Raza | Size-and shapedependent antibacterial studies of silver nanoparticles synthesized by wet chemical routes Nanomaterials[END_REF].

The appealing plasmonic properties and low optical losses of Ag NPs in the visible and nearinfrared spectral ranges have fueled intense research on the particles' localized surface plasmon resonance (LSPR) for the enhancement of their Raman scattering, fluorescence, and excitation, as well as the improvement of their photocatalysis, photoacoustic, and photovoltaic efficiency [START_REF] El-Fattah Z M, Mkhitaryan | Plasmonics in atomically thin crystalline silver films[END_REF][START_REF] Daoudi | Structural effects of silver-nanoprism-decorated Si nanowires on surface-enhanced Raman scattering[END_REF][START_REF] Futamata | Microscopic morphology and SERS activity of Ag colloidal particles Vibrational[END_REF][START_REF] Kalfagiannis | Plasmonic silver nanoparticles for improved organic solar cells Solar Energy Materials and[END_REF][START_REF] Petryayeva | Localized surface plasmon resonance: nanostructures, bioassays and biosensing--a review[END_REF][START_REF] Sujith | Surface enhanced Raman scattering analyses of individual silver nanoaggregates on living single yeast cell wall[END_REF]. In non-linear optical applications, metallic nanostructures are used to create "hot-spot" regions such that the electromagnetic energy from the incident light can be "squeezed" into the NPs' subwavelength spatial profiles. Such development in surfaceenhanced light-matter interactions and applications has motivated additional studies on the production of engineered Ag NPs with on-demand optical response [START_REF] Hu | A review of multiphysical fields induced phenomena and effects in spark plasma sintering: Fundamentals and applications Materials & Design[END_REF]. In addition to the nature and chemistry of the materials, the geometric morphology of the metallic nanostructures significantly affects their LSPR spectral properties and Surface-Enhanced Raman Scattering (SERS) intensities due to the corresponding changes in surface polarization [START_REF] Petryayeva | Localized surface plasmon resonance: nanostructures, bioassays and biosensing--a review[END_REF].

Indeed, the LSPR absorption of Ag NPs in the visible and infrared spectral regions may be tailored by varying the particle shape (sphere, triangle, cube, prism, bipyramid, octahedron, nanorod, nanoshell, nanostar, and structured array film) [START_REF] Lu | Chemical synthesis of novel plasmonic nanoparticles[END_REF][START_REF] Wiley | Maneuvering the surface plasmon resonance of silver nanostructures through shape-controlled synthesis[END_REF]. Local field enhancement for LSPR and SERS can also be found around sharp corners, edges, and conical tips [START_REF] Li | Mesoscopic and microscopic strategies for engineering plasmon-enhanced Raman scattering Advanced Optical Materials[END_REF].

A number of top-down and bottom-up methods of Ag nanostructure fabrication have been examined [START_REF] Bonafos | The 3D design of multifunctional silver nanoparticle assemblies embedded in dielectrics[END_REF][START_REF] Petryayeva | Localized surface plasmon resonance: nanostructures, bioassays and biosensing--a review[END_REF][START_REF] Sui | Plasmas for additive manufacturing Plasma Processes and[END_REF]. Among these methods, plasma-liquid electrochemistry can be used to rapidly synthesize colloidal Ag NPs by microplasma reduction of aqueous cations [START_REF] Kondeti | Ag + reduction and silver nanoparticle synthesis at the plasma-liquid interface by an RF driven atmospheric pressure plasma jet: Mechanisms and the effect of surfactant[END_REF][START_REF] Richmonds | Plasma-liquid electrochemistry: Rapid synthesis of colloidal metal nanoparticles by microplasma reduction of aqueous cations[END_REF]. After the deposition of Ag NPs by nebulization, inkjet printing, or spin-coating, the films can be modified by thermal annealing [START_REF] Abbas | Fabrication and characterization of silver thin films using physical vapor deposition, and the investigation of annealing effects on their structures[END_REF][START_REF] Moon | Thermal behavior of silver nanoparticles for low-temperature interconnect applications[END_REF][START_REF] Pandey | Effect of annealing temperature on morphological and optical transition of silver nanoparticles onc-plane sapphire[END_REF]. Depending on the conditions, various configurations and sizes of Ag NPs can be obtained, including irregular and round, dome-shaped nanostructures [START_REF] Pandey | Effect of annealing temperature on morphological and optical transition of silver nanoparticles onc-plane sapphire[END_REF]. However, the relatively high annealing temperatures are incompatible with a number of cost-effective polymer substrates such as polyethylene terephthalate, polyvinyl alcohol, and polyethylene, whose glass transition temperatures are below 150 o C [START_REF] Sui | Plasmas for additive manufacturing Plasma Processes and[END_REF].

The modifications induced by plasma processes performed under varying experimental conditions and configurations have also been assessed [START_REF] Hu | Low temperature sintering of Ag nanoparticles for flexible electronics packaging[END_REF][START_REF] Kang | Sintering of inkjet-printed silver nanoparticles at room temperature using intense pulsed light[END_REF][START_REF] Ma | Low temperature plasma sintering of silver nanoparticles[END_REF][START_REF] Sonawane | Effects of cold atmospheric plasma treatment on the morphological and optical properties of plasmonic silver nanoparticles[END_REF][START_REF] Srinivasan N R | Plasma treated activated carbon impregnated with silver nanoparticles for improved antibacterial effect in water disinfection[END_REF]. For instance, the efficiencies of scanning plasmas (e.g., microplasma jet) and large-area plasmas (e.g., lowpressure plasma and plane-to-plane dielectric barrier discharge) for spatially-localized treatments and spatially-homogeneous modification, respectively, have been investigated [START_REF] Sui | Plasmas for additive manufacturing Plasma Processes and[END_REF].

The very first study reported in the literature regarding plasma-modification of Ag NPs made use of a low-pressure Ar plasma with temperatures between 70 and 80 o C [START_REF] Reinhold | Argon plasma sintering of inkjet printed silver tracks on polymer substrates[END_REF]. In such systems, the plasma-generated species can effectively remove organics and promote surface diffusion, thereby sintering the Ag NPs into a dense network [START_REF] Hu | A review of multiphysical fields induced phenomena and effects in spark plasma sintering: Fundamentals and applications Materials & Design[END_REF][START_REF] Sui | Plasmas for additive manufacturing Plasma Processes and[END_REF]. Significant process optimization can be achieved through additional energy uptake; for example, through substrate heating by lowpressure plasma treatment or through the use of a warmer non-equilibrium plasma source at atmospheric pressure [START_REF] Wünscher | Towards single-pass plasma sintering: Temperature influence of atmospheric pressure plasma sintering of silver nanoparticle ink[END_REF]. For plasmonic applications, Branched Polyethylenimine (BPEI)functionalized Ag NPs with an average size of ~10 nm were recently exposed to a cold, atmospheric pressure plasma jet operated in argon and open to ambient air [START_REF] Sonawane | Effects of cold atmospheric plasma treatment on the morphological and optical properties of plasmonic silver nanoparticles[END_REF]. A reduction of the NP size was observed due to the removal of surface functional groups and subsequent surface oxidation. A blue shift of the LSPR absorption band and an increase in the SERS signal intensity of Rhodamine 6G at 612 cm -1 were also observed.

Inspired by plasma sintering studies realized at high temperatures [START_REF] Wünscher | Towards single-pass plasma sintering: Temperature influence of atmospheric pressure plasma sintering of silver nanoparticle ink[END_REF], this work explores the effect of a scanning open-to-air microwave argon plasma jet on the morphology and composition of thin films made of spin-coated Ag NPs (70-80 nm). Unlike the cold, atmospheric-pressure plasma jet treatment recently examined by Sonawane et al. [START_REF] Sonawane | Effects of cold atmospheric plasma treatment on the morphological and optical properties of plasmonic silver nanoparticles[END_REF], the microwave plasma jet used in this study is characterized by relatively high neutral gas temperatures in the range of 2000-2400 K [START_REF] Durocher-Jean | Characterization of a microwave argon plasma column at atmospheric pressure by optical emission and absorption spectroscopy coupled with collisional-radiative modelling[END_REF][START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF]. The effect of conventional thermal annealing (1300 K, 30 min) on the nanoparticles is also analyzed for comparison. Scanning electron microscopy, LSPR reflection spectroscopy, and X-ray photoelectron spectroscopy are used to examine the plasma-treated and annealed samples. Finally, hyperspectral Raman IMAging (RIMA) [START_REF] Gaufrès | Hyperspectral Raman imaging using Bragg tunable filters of graphene and other low-dimensional materials[END_REF][START_REF] Bigras | Probing plasma-treated graphene using hyperspectral Raman[END_REF] combined with multivariate curve resolution is used for Raman scattering analysis.

Experimental setup and diagnostics

Dispersions of Ag NPs (50 wt.%, d 50 = 70 nm, d 90 = 115 nm) in tripropylene glycol monomethyl ether were bought from Sigma Aldrich (796042-5G) and deposited as-is on thermal silica (500-nm-thick layer of SiO 2 on a 500 μm Si wafer) substrates by two-step spincoating. Before deposition, the uncoated substrates were cleaned by sequential ultrasonication in acetone, isopropanol, and water. In the first step of spin-coating, a homogeneous spread of the NP suspension was deposited onto the substrate during 30 s, at the spinning speed of 1000 rpm. In the second step, the spinner was run at 2000 rpm for 60 s to activate the drying process of the spin-coated layer. Afterwards, the substrate was annealed at 150 o C for 10 min to dry the coatings. The annealing conditions (time and temperature) were chosen such that they stimulate solvent evaporation without inducing significant agglomeration of Ag NPs [START_REF] Perelaer | Ink-jet printing and microwave sintering of conductive silver tracks[END_REF][START_REF] Reinhold | Argon plasma sintering of inkjet printed silver tracks on polymer substrates[END_REF].

Based on stylus profilometry measurements, the thickness of the spin-coated film prepared herein is around 300 nm.

The prepared samples of NP-coated SiO 2 /Si substrates were exposed to the jet of a microwave argon plasma. As illustrated in Figure 1a, the plasma jet was produced using a 2.45 GHz, 2 kW generator (Sairem) coupled to a WR430 surfaguide wave launcher [START_REF] Moisan | A waveguiide-based launcher to sustain long plasma columns through the propagation of an electromagnetic surface wave[END_REF], and the absorbed plasma power (incident minus reflected powers recorded by a bidirectional coupler connected to a bolometer) was fixed at 200 W. A capacitive coupler and a moveable plunger were used to bring the reflected power close to 0 W and to ensure maximum power transfer at the discharge tube location in the middle of the surfaguide. The plasma gas was introduced into the 83-mm-long fused silica discharge tube (26 mm above, 40 mm in, and 17 mm below the wave launcher; inner and outer diameters of 7 and 8 mm, respectively) from the top, at the mass flow rate of 2.1 SLM (standard liter per minute), then it expanded into the aluminum processing chamber. To simulate open-to-ambient air conditions, the aluminum chamber was fed with 25 SLM of clean, compressed air, and atmospheric pressure was maintained by gentle pumping. The flow of air was directed in such a way that it did not perturb the expansion of the plasma jet. Under such conditions, the plasma expanded beyond the end of the discharge tube [START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF][START_REF] Moon | Characteristics of an atmospheric microwave-induced plasma generated in ambient air by an argon discharge excited in an open-ended dielectric discharge tube[END_REF][START_REF] Redmond | Electrochemical Ostwald ripening of colloidal Ag particles on conductive substrates[END_REF] and thus, interacted with the grounded aluminum substrate holder (dimensions of 113 63 mm 2 ) placed in line with the tube outlet (see Figure 1b). In this study, the distance between the tube outlet and the substrate surface was fixed at 10 mm. A set of two Velmex Inc. stepping motors were used for two-dimensional plasma treatment of spin-coated Ag NPs. The substrate holder was scanned along one axis at the speed of 50 mm/s, with a step size of 5 mm along the 2 nd axis. Each plasma pass was programmed to cover an area of 85 50 mm 2 surface. Considering that the size of a microwave plasma jet is typically ~5 x 5 mm 2 and that the substrate surface area is 15 15 mm 2 , the in-plasma ( i.e. time during which a domain of the sample surface is exposed to the argon plasma jet in each pass) and outof-plasma (i.e. time during which the domain is not exposed to the plasma in each passes, excluding the time between consecutive passes) times are roughly ~1 s and ~16 s, respectively.

Therefore, Ag NPs were not continuously exposed to the plasma. Knowing that ~25 s are needed for the plasma to go back to its initial position, the total out-of-plasma time during each pass is more than 41 s.

Optical emission spectroscopy measurements were carried out at 1 mm above the substrate holder using a high-resolution spectrometer (HORIBA Jobin Yvon THR1000 monochromator) equipped with a photomultiplier tube (Hamamatsu R636-10) and a holographic grating of 1800 lines/mm. This system was used to measure the line broadening of the Ar 2p 2 -1s 2 and 2p 3 -1s 2 (Paschen notation) emission lines at 826 and 841 nm, respectively.

These two lines are strongly affected by resonance broadening in atmospheric pressure plasmas, and the Lorentzian portion of their broadening is completely dominated by mechanisms related to the neutral gas temperature ( ) (resonance and van der Waals broadening) [START_REF] Durocher-Jean | Characterization of a microwave argon plasma column at atmospheric pressure by optical emission and absorption spectroscopy coupled with collisional-radiative modelling[END_REF][START_REF] Labelle | On the rotational-translational equilibrium in non-thermal argon plasmas at atmospheric pressure Plasma Sources[END_REF]. A typical measurement along with the corresponding Voigt fit is presented in Figure 2. Based on detailed line broadening analysis, the experimental conditions investigated herein lead to a value in the order of 2200 ± 200 K. This value agrees well with those reported by Durocher-Jean et al. [START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF] for microwave plasma jets; however, it is considerably higher than the values commonly observed for low and mid-frequency plasma jets ( ~300-550 K [START_REF] Hofmann | Power dissipation, gas temperatures and electron densities of cold atmospheric pressure helium and argon RF plasma jets Plasma Sources[END_REF][START_REF] Sainct | Spatially-resolved spectroscopic diagnostics of a miniature RF atmospheric pressure plasma jet in argon open to ambient air[END_REF]). Considering that the population of the Ar 1s state is sufficiently small within the range of experimental conditions employed herein (typically below 10 9 cm -3 [START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF]), the radiation trapping of Ar 2p-to-1s transitions was neglected [START_REF] Durocher-Jean | Characterization of a microwave argon plasma column at atmospheric pressure by optical emission and absorption spectroscopy coupled with collisional-radiative modelling[END_REF].

Morphological analysis was conducted using a JEOL 7600 field emission scanning electron microscope (FE-SEM). Prior to each analysis, the samples were coated with a thin layer of conductive gold to improve the electrical conductivity across the whole sample surface. The 9 500 ms). To determine whether Ag oxidation occurs only at the nanoparticle surface or more profoundly in the nanoparticle volume, X-ray diffraction analysis of the plasma-treated and annealed samples was also conducted.

The LSPR properties of the prepared samples were analyzed by reflection spectroscopy using a halogen light source and a low-resolution UV-Vis-NIR spectrometer, at normal incidence between 350 and 800 nm. Meanwhile, the Raman signature was assessed by hyperspectral Raman spectroscopy using the RIMA (Raman IMAger) system [START_REF] Gaufrès | Hyperspectral Raman imaging using Bragg tunable filters of graphene and other low-dimensional materials[END_REF][START_REF] Bigras | Probing plasma-treated graphene using hyperspectral Raman[END_REF], which provides uniform laser illumination ( laser = 532 nm) over a wide area (130 × 130 m 2 ) while maintaining high selectivity and diffraction-limited spatial resolution (~390 nm). In this work, the laser power was set to 750 mW (power density of 4.4 x 10 7 W/m 2 ).

Results and discussion

Figure 3a depicts a magnified SEM image of spin-coated Ag NPs on thermal silica substrate.

Good coverage of the substrate is observed, with few randomly dispersed inhomogeneities on the surface. This confirms that the surface wetting of the substrate during the spin coating procedure is good. Furthermore, the Ag NPs are mostly rounded, and only a few have a polygonal shape. To determine the Feret diameter distribution of as-deposited NPs, a full set of SEM images obtained at the same magnification were processed using the ImageJ software [START_REF] Papadopulos | Common tasks in microscopic and ultrastructural image analysis using[END_REF]. The results shown in Figure 3b demonstrate that the NPs of untreated samples exhibit an asymmetric size distribution that is centered around 70-80 nm, which is consistent with the values provided by the supplier (d 50 = 70 nm). Overall, the SEM data reveal that the Ag NPs do not agglomerate during the two annealing steps carried out at 150 o C for 10 min [START_REF] Perelaer | Ink-jet printing and microwave sintering of conductive silver tracks[END_REF][START_REF] Reinhold | Argon plasma sintering of inkjet printed silver tracks on polymer substrates[END_REF]. Ag NPs (Figure 3a) to a complex, multi-level film microstructure (Figures 4a and4c). As shown in Figure 4e, neighboring particles at the bottom of the features in plasma-treated samples (100 passes) experience both, neck formation and particle coarsening, similar to Ag NPs sintered by low-temperature plasmas [START_REF] Ma | Low temperature plasma sintering of silver nanoparticles[END_REF]. Considering that particle coarsening is driven by polydispersity [START_REF] Shi | Relations between coarsening and densification and mass transport path in solid-state sintering of ceramics: Model analysis[END_REF], the morphological change induced by plasma treatment may be attributed to the association of isolated Ag NPs, a phenomenon that is often observed after thermal-or plasmaassisted decomposition of the surrounding organic shells [START_REF] Mitronika | Modification of the optical properties and nano-crystallinity of anatase TiO 2 nanoparticles thin film using low pressure O 2 plasma treatment Thin Solid Films[END_REF][START_REF] Sui | Plasmas for additive manufacturing Plasma Processes and[END_REF]. After 100 passes, polygonal microstructures with very sharp edges are observed (Figures 4c and4d). Although the sidewalls of these microstructures appear to be smooth, 50-100 nm features remain on the topmost surface (Figure 4f). This suggests that particle growth proceeds according to the so-called Ostwald ripening mechanism: larger particles grow at the expense of smaller ones [START_REF] Gentry | Ostwald ripening in metallic nanoparticles: Stochastic kinetics The[END_REF][START_REF] Redmond | Electrochemical Ostwald ripening of colloidal Ag particles on conductive substrates[END_REF][START_REF] Zong | Kinetically controlled seed-mediated growth of narrow dispersed silver nanoparticles up to 120 nm: secondary nucleation, size focusing, and Ostwald ripening[END_REF].

Such ripening occurs because large particles are more energetically favored than the small ones. 13

The SEM images recorded over various scales were used to determine the particle size distribution of spin-coated Ag NPs after plasma treatment. The obtained results (Figure 5) show that the average size of Ag NPs drastically increases from 70-80 nm to 1-10 µm after exposure to the microwave argon plasma jet. The relative concentration of large particles increases slightly upon increasing the number of jet passes from 10 to 100. As detailed above, the in-plasma time during each pass is approximately 1 s. Therefore, each domain in the sample surfaces treated with 10 and 100 passes is exposed to the plasma jet for ~10 and ~100 s, respectively, with several tens of seconds between two consecutive passes. Hence, the "average" surface temperature of treated samples is lower than the "transient" temperature attained during the period of exposure to the plasma jet (neutral gas temperature = 2200 ± 200 K based on the data in Figure 2). The variation in surface temperature (T s ) was continuously monitored using a thermocouple. To prevent direct plasma exposure and damage, the thermocouple was placed inside the substrate holder at 3 mm below the topmost surface (see Figure 1). Since the holder is made of aluminum, the temperature measured below the substrate is expected to be roughly the same as that attained at the topmost surface (T s ). As shown in Figure 6, T s , initially at 300 K, increases significantly within the first few minutes of plasma-on time. Note that each oscillation in the figure represents a single two-dimensional scan of the substrate holder (85 50 mm 2 surface, with the substrate surface covering only 15 15 mm 2 ). T s continues to rise after 10 passes, and saturation is reached at about 50 two-dimensional scans. Based on the data presented in Figure 6, the average surface temperature of samples treated with 100 passes of plasma is 450 K, which cannot promote significant atomic diffusion and the formation of micrometer-scale particles over short plasma treatment times. Therefore, other plasma-surface interactions involving temporally localized energy transfer processes must be occurring during plasma treatment. These interactions include surface bombardment by high-energy ions and very-hot neutral species, heterogeneous surface neutralization of positive ions, and heterogeneous surface de-excitation of metastable and vibrational states, among others [START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF][START_REF] Ma | Low temperature plasma sintering of silver nanoparticles[END_REF], all of which may contribute to and induce the complex morphological evolution observed herein. Considering that the experimental conditions applied in this study lead to a low kinetic energy of positive ions, the interaction of these ions with the substrate surface is expected to be minimal, and thus, ion-induced energy transfer is mostly governed by heterogeneous surface neutralization of positive ions (transfer of 15.76 eV energy for Ar + ). In addition, the microwave energy is entirely coupled to the atmospheric-pressure plasma jet such that microwave-assisted heating or electric-field-induced phenomena can reasonably be ignored [START_REF] Perelaer | Ink-jet printing and microwave sintering of conductive silver tracks[END_REF].Given that the number densities of charged species near the substrate surface are very low, no significant charging effects are expected [START_REF] Darny | Analysis of conductive target influence in plasma jet experiments through helium metastable and electric field measurements[END_REF][START_REF] Goldberg | 1D time evolving electric field profile measurements with sub-ns resolution using the E-FISH method[END_REF][START_REF] Xiong | Atmospheric pressure ionization waves propagating through a flexible high aspect ratio capillary channel and impinging upon a target Plasma Sources[END_REF]. To further assess the sintering effect of the microwave plasma jet, the morphology of plasma-treated samples was compared to that of thermally annealed spin-coated Ag NPs. The annealing process was conducted in a partially sealed oven that was continuously filled with argon gas. The oven temperature was varied as per the three-ramp temperature profile illustrated in Figure 7a, and representative SEM images recorded after each ramp are presented in Figures 7b-7d. Based on these figures, the films annealed at low temperatures exhibit a dramatic change in morphology from discrete and spherical Ag NPs (Figure 3a) to a continuous and sintered microstructure (Figure 7b). Such behavior is typically reported for thermallyactivated sintering [START_REF] Abbas | Fabrication and characterization of silver thin films using physical vapor deposition, and the investigation of annealing effects on their structures[END_REF][START_REF] Ahn | Omnidirectional printing of flexible, stretchable, and spanning silver microelectrodes[END_REF][START_REF] Moon | Thermal behavior of silver nanoparticles for low-temperature interconnect applications[END_REF][START_REF] Pandey | Effect of annealing temperature on morphological and optical transition of silver nanoparticles onc-plane sapphire[END_REF][START_REF] Shen | Preparation of solid silver nanoparticles for inkjet printed flexible electronics with high conductivity[END_REF]. Thermally-enhanced atomic diffusion continues at higher annealing temperatures, resulting in the formation of micrometer-scale spherical structures, as shown in the SEM images recorded after the 2 nd and 3 rd temperature ramps (Figures 7c and7d, respectively). Again, heat-assisted growth occurs because larger particles are more energetically favored than smaller ones [START_REF] Gentry | Ostwald ripening in metallic nanoparticles: Stochastic kinetics The[END_REF][START_REF] Redmond | Electrochemical Ostwald ripening of colloidal Ag particles on conductive substrates[END_REF][START_REF] Zong | Kinetically controlled seed-mediated growth of narrow dispersed silver nanoparticles up to 120 nm: secondary nucleation, size focusing, and Ostwald ripening[END_REF]. Detailed analysis of many SEM images shows that the particle size of thermally annealed Ag NPs varies in the range of 1-4 µm (Figure 5). Notably, the structures obtained after thermal annealing are round, whereas those obtained after plasma treatment are polygonal (sharp edges and corners). Compared to the temporally localized energy uptake occurring during the latter process, the former process exhibits a continuous energy uptake, which promotes surface diffusion phenomena and increases the surface layer continuity. Based on previous studies, surface diffusion is the main driving force behind Ag particle evolution at temperatures below 850 K [START_REF] Pandey | Effect of annealing temperature on morphological and optical transition of silver nanoparticles onc-plane sapphire[END_REF]. Here, it must be noted that even though the melting temperature of bulk silver substrates is 1262 K, the temperature needed for sintering of Ag NPs is reduced to 450-500 K due to the well-known Gibbs-Thomson effect [START_REF] Alshehri | Enhanced electrical conductivity of silver nanoparticles for high frequency electronic applications[END_REF][START_REF] Moon | Thermal behavior of silver nanoparticles for low-temperature interconnect applications[END_REF].

Plasma-and thermally-induced modifications of spin-coated Ag NPs were also analyzed by XPS to examine possible surface oxidation. Knowing that the reported binding energy values determined based on the most intense Ag 3d peak exhibit a large discrepancy, the Ag MNN Auger structure was analyzed instead. As highlighted by several authors [START_REF] Kaushik | XPS core Liwel spectra and Auger parameters for some silver compounds[END_REF][START_REF] Lundwall | Enhanced surface sensitivity in AES relative to XPS observed in free argon clusters Surface[END_REF], the chemical shift of this Auger system is larger than that of the photoelectron counterparts and is independent of sample charge shifts [START_REF] Ferraria | X-ray photoelectron spectroscopy: Silver salts revisited[END_REF]. The XPS spectra shown in Figure 8 The oxidation of spin-coated Ag NPs was confirmed by EDS mapping. The results illustrated in Figure 9 demonstrate that after exposure to 100 passes of the microwave argon plasma jet, oxygen is detected at the bottom of the silver features, as well as on top of the plasmagenerated microstructures. The Si signal attributed to the substrate is also detected; however, its intensity is lower at the locations of the micrometer-scale features than at other locations. This is consistent with a significant augmentation in surface topology after plasma treatment (see SEM images in Figure 4). As expected, the opposite is observed for the silver signal, with a more intense Ag signal detected on the microstructures. XRD spectroscopy analyses further confirm that plasma treatment results in the oxidation of Ag NPs. Indeed, the spectra presented in Figure 10 demonstrate that peaks corresponding to Ag 2 O appear at 2θ of 33 o and 55 o only after treatment with microwave plasma jet. The intensities of these peaks increase with increasing number of jet passes. In contrast, the thermally annealed sample exhibits no Ag 2 O signals in its XRD spectrum, and only signals corresponding to metallic silver are detected (Figure 10). Hence, despite the surface oxidation evidenced by XPS for both, plasma and annealing treatments, the XRD data reveal that only the former leads to volume oxidation of Ag NPs.

20

The LSPR signal of Ag NPs deposited on SiO 2 /Si substrates was examined using reflection spectroscopy, and the variation of reflectivity R (in arbitrary units) as a function of wavelength  is presented in Figure 11. R() is defined as the ratio of the emission intensity reflected at normal incidence I() to that of the lamp light reflected on a perfectly reflective mirror placed at the substrate location I 0 (). Before plasma treatment, R() exhibits little variation, with only a small drop linked to LSPR extinction around 400 nm. This agrees well with LSPR calculations based on the discrete dipole approximation for Ag NPs [START_REF] Lu | Chemical synthesis of novel plasmonic nanoparticles[END_REF]. After 10 passes of the microwave argon plasma jet, a redshift in the main extinction band is observed, which is consistent with the formation of larger structures. The reflection spectrum does not change significantly upon increasing the number of passes from 10 to 100; however, the extinction signal becomes broader. Since the samples treated with 10 and 100 passes exhibit comparable size distributions (see Figure 5), the broader extinction signal observed after 100 passes is most likely due to the formation of a greater number of polygonal-like microstructures characterized by sharp corners and edges (see Figures 4c and4d) [START_REF] Li | Mesoscopic and microscopic strategies for engineering plasmon-enhanced Raman scattering Advanced Optical Materials[END_REF]. Another dip in reflectance is detected around 640 nm, particularly after 100 passes of the plasma jet. This signifies an increase in absorbance, and thus, in LSPR. A redshift in the extinction signal is also observed after thermal annealing (ramp 3), and it is attributed to the formation of larger microstructures. No other extinction is detected, which agrees well with the much rounder edges of Ag NPs displayed in Figure 7.

The Raman signals of the samples were also recorded using the RIMA system [START_REF] Gaufrès | Hyperspectral Raman imaging using Bragg tunable filters of graphene and other low-dimensional materials[END_REF][START_REF] Bigras | Probing plasma-treated graphene using hyperspectral Raman[END_REF], and spatially-averaged Raman spectra were collected over a 29 x 29 µm 2 area. As shown in Figure 12, the Si/SiO 2 substrate spectrum exhibits a continuously decreasing signal in the wavenumber range of 850-1000 cm -1 . The same signal is observed in the spectrum of the untreated sample (spin-coated Ag NPs on the substrate); however, its intensity increases with increasing wavenumber, and small peaks appear in the same spectral region. Broad bands centered at 1340 and 1590 cm -1 are also detected. According to previous studies, these bands may be attributed to the SERS signals of carbon species surrounding the Ag NPs [START_REF] Li | Ag@C core-shell colloidal nanoparticles prepared by the hydrothermal route and the low temperature heatingstirring method and their application in surface enhanced Raman scattering[END_REF][START_REF] Moyer | Surface-enhanced Raman scattering spectroscopy of single carbon domains on individual Ag nanoparticles on a 25 ms time scale[END_REF]. After treatment with 10 passes of the microwave plasma jet, the carbon-related signals become less intense, and they disappear completely after 100 passes. This is consistent with the hypothesis of plasma-assisted decomposition of the organic shell surrounding the spin-coated Ag NPs [START_REF] Mitronika | Modification of the optical properties and nano-crystallinity of anatase TiO 2 nanoparticles thin film using low pressure O 2 plasma treatment Thin Solid Films[END_REF]Figure 11. Reflection spectrum at normal incidence of the spin-coated Ag NPs before and after treatment with 10 plasma passes, 100 plasma passes, and thermal annealing (ramp 3).

22 60]. In addition, a new single peak is observed at around 960 cm -1 after plasma treatment, and its intensity increases with increasing number of passes. Notably, equal Raman shifts of this peak are detected at the laser excitation wavelengths of 532 (Figure 12) and 488 nm (not shown), which confirms that it is due to Raman scattering and not to fluorescence excitation.

The 960 cm -1 peak has also been observed for polycristalline silver exposed to oxygen at temperatures between 25 and 500 o C [START_REF] Wang | Interaction of polycrystalline silver with oxygen, water, carbon dioxide, ethylene, and methanol: In situ Raman and catalytic studies[END_REF] (the high temperatures in this range are similar to that of the warm plasma employed herein). Like the plasma-treated samples, the annealed sample does not exhibit carbon-related signals in its Raman spectrum; however, the 960 cm -1 peak is also not observed. This is consistent with the XRD results showing volume oxidation of Ag NPs by plasma treatment only.

To confirm the origin of the surface-enhanced Raman scattering observed in Figure 12, spatially-resolved Raman spectra were recorded using the RIMA system. To the best of our knowledge, this work represents the first application of RIMA to Ag-based nanostructures and Considering that minimal power density was applied in this study to minimize surface degradation by laser irradiation, the number of obtained data sets is low. To overcome this limitation and accurately analyze the spectral signatures, the Raman spectra were processed by multivariate curve resolution (MCR), according to the log-likelihood maximization (LLM) method [START_REF] Fauteux-Lefebvre | A hierarchical multivariate curve resolution methodology to identify and map compounds in spectral images[END_REF]. This method is generally used to identify and extract typical spectral signatures or components, allowing to account for the whole data set and to obtain the relative contribution maps of different components from a given spectral data set. In this framework, the decomposition can be expressed as X(M N) = C (M K) S T (K N) + E(M N), where X is the raw data matrix, S T is the component, C is the relative contribution of the component, and E is the residuals. M, N , and K are the pixel number, the number of wavenumber energy levels, and the number of components, respectively. Within the range of experimental conditions investigated herein, three distinct spectral signatures are determined by MCR, resulting in three relative spectra, as shown on the bottom of Figure 13. The first spectrum (S0) exhibits low-density regressive behaviour in the 850-1050 cm -1 spectral range, with a small artifact feature [START_REF] Gaufrès | Hyperspectral Raman imaging using Bragg tunable filters of graphene and other low-dimensional materials[END_REF][START_REF] Bigras | Probing plasma-treated graphene using hyperspectral Raman[END_REF].

Meanwhile, the second spectrum (S1) shows a prominent peak with high amplitude near 960 cm -1 . Finally, the third spectrum (S2) depicts a signal with a steady intermediate intensity that slightly increases with increasing wavenumber. By capitalizing on the RIMA and MCR capabilities, spatially-resolved values of the relative contributions C of each component (summation of the contributions of the the three components is equal to 1) were obtained (Figure 13).

Based on the results presented in Figures 12 and13, S0, S1, and S2 can be linked to the spectral signatures of specific features. The first component S0 represents a vast majority of the Raman signal from the Si/SiO 2 substrate, as shown in Figure 13, whereas the third component S2, mostly accounts for the signal of the untreated sample and thus, is related to the SERS signals of carbon species surrounding the Ag NPs [START_REF] Li | Ag@C core-shell colloidal nanoparticles prepared by the hydrothermal route and the low temperature heatingstirring method and their application in surface enhanced Raman scattering[END_REF][START_REF] Moyer | Surface-enhanced Raman scattering spectroscopy of single carbon domains on individual Ag nanoparticles on a 25 ms time scale[END_REF]. After exposure to 100 passes of the microwave argon plasma jet, the Ag NPs on the sample surface are oxidized, and the carbon- dominated Raman signature shifts to a strong Raman signal of oxidized silver near 960 cm -1 (i.e.

from mostly S2 to mostly S1). By comparing the optical images with S1 and S2 mappings, it can be seen that the strong Raman signal provided by S1 is attributed to large-scale domains that are dominated by micrometer-scale polygonal structures with sharp edges and corners. Hence, in the absence of the local-field-enhancement effects observed after thermal annealing of spincoated Ag NPs, the contribution from S1 vanishes.

Conclusion

In

Figure 1 .

 1 Figure 1. (a) Schematic representation of the microwave plasma jet open to ambient air. (b) Picture of the plasma jet interacting with the substrate holder (distance between the tube outlet and the top surface = 10 mm). The red square indicates the position of the thermocouple (3 mm below the topmost surface) used for surface temperature measurements.

  acceleration voltage was set to 15 kV for untreated samples and to 5 kV for plasma-treated and annealed samples (details on annealing conditions provided below). The working distance ranged between 3.7 and 12.8 mm, depending on the sample. Selected samples were also probed by energy-dispersive X-ray spectroscopy (EDS) to acquire a mapping of the surface chemistry. The surface roughness and topology of these samples were analyzed by atomic force microscopy (MultiMode 8-Mode Scan from Bruker) using a standard etched-silicon cantilever operated in peak force quantitative nanomechanical tapping mode. The Nanoscope Analysis software v1.5 (Bruker) was used to analyze the recorded images (1.8 1.8 µm 2 ). To determine the chemical composition of the topmost surface of samples, X-ray photoelectron spectroscopy (XPS) analysis was performed at 1486.6 eV using a Kratos Nova spectrometer equipped with Al Kα radiation. Based on the work of Ferraria et al.[START_REF] Ferraria | X-ray photoelectron spectroscopy: Silver salts revisited[END_REF], the Ag MNN Auger structure of spincoated Ag NPs was recorded at the pass energy of 40 eV (step scan = 0.1 eV and dwell time =

Figure 2 .

 2 Figure 2. High resolution emission spectrum of the Ar 2p 3 -1s 2 transition at 826 nm and the Ar 2p 8 -1s 4 transition at 841 nm. The Voigt, Gaussian, and Lorentzian broadenings for each line are also presented.

Figure 4

 4 Figure 4 compares the SEM images (low-and high-magnification) of spin-coated Ag NPs before and after treatment with multiple passes ( 10 or 100) of the microwave argon plasma jet. Clearly, plasma treatment transforms the morphology of the samples from isolated, round

Figure 3 .

 3 Figure 3. (a) SEM analysis of silver nanoparticles deposited on SiO 2 /Si substrate by spin-coating. NPs with different diameters are highlighted by red circles. (b) Particle size distribution of the spin-coated silver nanoparticles obtained by analysis of the SEM and AFM images of multiple samples.

Figure 4 .

 4 Figure 4. SEM images of silver nanoparticles (at different magnifications) after treatment with (a,b) 10 and (c-f) 100 passes of a microwave argon plasma jet. (e) and (f) show magnified images of the micro-sized structures and the region between them, respectively.

Figure 5 .

 5 Figure 5. Particle size distributions of the spin-coated Ag Nps treated with 10 and 100 passes of microwave plasma jet, as well as the particles subjected to thermal annealing (ramp 3, see details below).

Figure 6 .

 6 Figure 6. Variation of surface temperature (recorded using a thermocouple placed under the substrate holder) as a function of total plasma-on time. Here, each oscillation represents a single two-dimensional scan of the substrate holder.

Figure 7 .

 7 Figure 7. (a) Scheme of the three-ramp temperature profile used during annealing treatments. (b-d) SEM images of Ag NPs recorded after (b) ramp 1, (c) ramp 2, and (d) ramp 3.

  demonstrate that all samples exhibit two peaks corresponding to Ag M 5 N 45 N 45 and Ag M 4 N 45 N 45 Auger electrons in the Ag MNN region. In the untreated sample spectrum, these peaks are detected at 353 and 358 eV; however, they are shifted by about 2.5 eV (to lower kinetic energies) after exposure to 10 passes of the open-to-air microwave argon plasma jet. Similar chemical shifts and peak shapes are observed for samples treated with 100 plasma passes or thermal annealing. According to Ferraria et al. [16], the chemical shifts of the Ag M 4 N 45 N 45 peaks of AgO and Ag 2 O are 2.1 and 2.6 eV, respectively, and similar values are reported for the Ag M 5 N 45 N 45 peak. Therefore, surface oxidation of spin-coated Ag NPs by plasma treatment or thermal annealing leads to Ag 2 O as the dominant oxidation state. Considering that the Auger structures observed after 10 and 100 plasma jet passes are similar, the surface oxidation is relatively fast.

Figure 8 .

 8 Figure 8. Ag MNN spectrum of the spin-coated Ag NPs before and after treatment with10 jet passes, 100 jet passes, or thermal annealing (ramp 3).

Figure 9 .

 9 Figure 9. EDS analysis of spin-coated Ag NPs treated with 100 passes of the microwave argon plasma jet. The results are shown for oxygen, silicon, and silver.

Figure 10 .

 10 Figure 10. XRD spectra of spin-coated Ag NPs before and after treatment with the microwave argon plasma jet (10 and 100 passes) and thermal annealing.

Figure 12 .

 12 Figure 12. RIMA spectra of spin-coated Ag NPs before and after treatment with the plasma jet (10 or 100 passes) or thermal annealing (ramp 3). The Raman spectrum of the SiO 2 /Si substrate is also presented for comparison.

Figure 13 .

 13 Figure[START_REF] Durocher-Jean | Deposition of anti-fog coatings on glass substrates using the jet of an open-to-air microwave argon plasma at atmospheric pressure Plasma Processes and[END_REF]. Deconvolution of spatially-resolved RIMA images processed by multivariate curve resolution, according to the log-likelihood maximization method. Mappings of the relative contributions C of the three different components obtained (S0, S1, and S2) for the substrate, the untreated sample, the plasma-processed sample, and the thermally-annealed sample.Optical images recorded over the same domain are also shown for comparison.

  summary, this work examines the modification of thin films made of spin-coated silver nanoparticles by a warm microwave argon plasma jet open to ambient air. Scanning electron microscopy analysis shows that plasma jet treatment results in the formation of polygonal microstructures with sharp edges and corners, unlike the spherical microstructures obtained with conventional thermal annealing. Both treatments (plasma and thermal annealing) lead to surface oxidation, as evidenced by X-ray photoelectron spectroscopy and energy-dispersive X-ray spectroscopy analyses. Interestingly, a rise in localized surface plasmon resonance and surface-enhanced Raman scattering signals is observed for plasma-treated samples but not for the annealed ones. Multivariate curve resolution processing of the Raman mappings reveals that Raman signal is controlled by large-scale, polygonal, and sharp-edged domains. Although previous studies highlight the strong potential of low-temperature plasma jet treatment for open-to-ambient-air plasma sintering or blue-shift LSPR tuning of Ag NPs, this work shows that additional SERS enhancement with red-shift LSPR tuning may be achieved by exposing Ag nanoparticles to a warmer microwave argon plasma jet.

Acknowledgments

The equipment used in this work was acquired through the Leaders Opportunity Fund of the Canadian Foundation for Innovation (CFI). Financial supports from the National Science and Engineering Research Council (NSERC), the Canada Research Chair program, the Directions des Relations Internationales of the Université de Montréal, and the CNRS are acknowledged. This work was realized in the context of the International Research Network on Nanomatériaux Multifonctionnels Contrôlés (IRN-NMC).