Antoine Rouilly

Luc Rigal

Robert G Gilbert
email: gilbert@chem.usyd.edu.au

Synthesis and properties of composites of starch and chemically modified natural rubber

Keywords: Starch, Natural rubber latex, Polymer colloid

A means is developed for forming polysaccharide-based composites with useful material properties through use of unmodified and chemically modified natural rubber latex (NRL). Starch was used as a model for polysaccharides. The NRL was modified by grafting with dimethylaminoethyl methacrylate (DMAEMA) to form a latex with cationic water-soluble polymeric 'hairs' of polyDMAEMA, which should form hydrogen bonds with starch. Starch solutions, containing 20% glycerol as a film-forming aid, and the modified NRL were mixed and films allowed to form. The unmodified latex acted only as filler in the starch films, but with modified NRL, the mechanical properties of the films were significantly altered. The elastic modulus was greatly decreased and strain at break greatly increased. The glass transition temperature increased from K48 8C to K32 8C, suggesting significant compatibilization. Freeze-fracture TEM micrographs indicate strong interactions between the surface of the modified NRL and starch. The polyDMAEMA chains are more hydrophilic than the starch, and the addition of grafted latex results in a 208 drop of the water contact angle of the formed film, and a 25% increase of the water absorption compared to the native starch; with unmodified NRL, the opposite effect was observed.

Introduction

The use of polymers from renewable resources is an environmentally advantageous alternative to synthetic polymers in some applications. Many polysaccharide biopolymers have been examined in this context [1], with the long-term aim of value-adding to 'waste' agricultural byproducts such as sugarbeet pulp and rice husks. Pure starch is a good model for such biopolymers, and also is useful as a substrate in its own right.

The thermoplastic properties of starch have been extensively studied [START_REF] Tomka | Water relationships in food[END_REF], and are directly related to its water content; in its dry state, the glass transition temperature (T g ) of starch is above its degradation temperature. While many technologies have been used to process starch, the easiest way of obtaining a film is by casting from a solution. To obtain useful materials from starch, the native properties must be enhanced, because of starch's high water sensitivity and poor mechanical properties compared to those of synthetic polymers. The influence of water content [START_REF] Hulleman | [END_REF] and of external plasticizers like glycerol [START_REF] Tomka | The glassy state in foods[END_REF] and sorbitol [START_REF] Gaudin | [END_REF] have been investigated to decrease the brittleness of these materials. Vegetable fibres [6] and mineral fillers [7,8] have been used to enhance the mechanical strength of the starch. Starch and related biopolymers have been chemically modified [9,10] or blended with synthetic polymers [11,12] to improve mechanical strength and water resistance.

As a suitable polymer for use in starch modification, natural rubber latex (NRL) has many attractions. It is a latex, which facilitates blending with a starch solution and indeed with any particulate substance such as sawdust or pulp; it is a renewable resource, can be biodegraded [13], and contains natural stabilizers (i.e. proteins and lipids) which should help compatibilization with starch. The polymer is essentially 100% cis-1,4-polyisoprene, and this conformation leads to a number of useful mechanical properties such as improved mechanical strength on stretching. It is also inexpensive. NRL is supplied as the natural latex (which has a broad particle size distribution, ranging from 100 nm to 2 mm), stabilized with ammonia. The abbreviation NRL will always be used for this 'highammonia' latex.

Recently, a method has been developed to modify NRL by grafting a 'hairy layer' of hydrophilic polymers [14]. Of the various hydrophilic monomers which could be used for this purpose, that chosen here is dimethylaminoethyl methacrylate, DMAEMA. This grafting procedure has been shown by NMR [15] and colloidal stability studies [14,16] to yield covalently bonded polyDMAEMA hairs as well as some free polyDMAEMA in the water phase [16]. The amine functionality of polyDMAEMA confers considerably enhanced colloidal stability on NRL: while NRL coagulates when the pH is taken below w8.5, the polyDMAEMA-modified NRL is stable to pH values as low as 2 [14]. DMAEMA is biodegradable, although has some ecotoxicity [17].

The aim of this work is to use NRL and NRL modified by polyDMAEMA directly in the wet casting of starch films and to investigate the effect on the properties of these films, with the long-term goal of value-adding to agricultural 'waste' byproducts. Alternatives approaches on starch/rubber blends have examined the use of starch as a filler in a rubber matrix [18,19] with phase compatibilization by chemical modification of either the rubber [20][21][22][23] or the starch [24][25][26].

Materials and methods

Materials

Rice starch (chosen because this is an important product in Australia) was supplied by Sigma. Because it has been observed [18] that the amylose content of starch can affect mechanical properties in starch/rubber blends, the amylose content of this sample was measured as follows. Starchiodine complexes were made for the starch following the standard method for quantifying amylose [27,28], and absorbance was measured every 1.25 nm over the visible spectrum using a scanning spectrophotometer (GBC UV/ VIS 918). The amylose content was thus found to be 28%. The present data are for this fixed amylose content. NRL ('high ammonia latex', 61% solids) from RLA Polymers was used as supplied. Dimethylaminoethyl methacrylate (DMAEMA), cumene hydroperoxide (CHP), tetraethylenepentamine (TEPA), all Aldrich Reagent Grade, were used as received. Reagent grade glycerol and analytical grade Mg(NO 3 ) 2 , K 2 SO 4 , dimethyl sulfoxide (DMSO) and acetone were used as received.

Grafting procedure

To graft DMAEMA onto NRL, the following procedure was used, following that of Lamb et al. [14]. The weight of DMAEMA was chosen to be 5, 10 and 15% w/w with respect to dried weight of rubber in the NRL, these grafted rubber latexes being denoted GRL5, GRL10 and GRL15. NRL, DMAEMA, CHP, TEPA and 2.5% ammonia solution were combined in the amounts given in Table 1. The reaction mixture was agitated with a low-shear impeller, a 10% aqueous solution of TEPA was then added shot-wise over 1 h at 5 min intervals, and the system was cooled in an ice/water bath during the first 8 h of reaction before gradually warming to room temperature, and allowed to continue react for another 16 h. After reaction, the latex was filtered through glass wool and neutralized to pH 7 with 1 M hydrochloric acid solution. The resulting latex as used in the present paper contains both grafted and ungrafted poly-DMAEMA [16]. For reasons detailed elsewhere [16], the amount of grafting is extremely hard to quantify, because of the difficulty of quantitatively removing all ungrafted polyDMAEMA. However, NMR studies on the gel fraction [15], which will contain the lowest amount of grafted polymer [15], show that the lower bound for the amount of grafted polyDMAEMA in the present system is 5 mol%.

Film formation

Granular starch was dispersed in ultra-pure water (3% w/w) and gelatinized by stirring and heating to 75 8C followed by dissolution in an autoclave at 120 8C under 120 kPa pressure for 30 min. A clear, viscous solution was obtained. To this solution, glycerol (as plasticizer for the starch), 20% w/w relative to starch on a dry basis, was added. Starch/latex samples with ratios of 100/0, 95/05, 90/10, 85/15, 80/20 and 70/30 (w/w on a dry basis, that of the latex referring to the solid content obtained after drying at 60 8C) were made by adding latex to the starch/glycerol solution. The system was homogenized by vigorous shaking of the mixture. This was then poured on polystyrene dishes and allowed to dry at 25 8C over 3 days. The resulting films (20-50 mm in thickness) were then dried at 60 8C overnight and maintained at 54% relative humidity (r.h.) over a saturated solution of Mg(NO 3 ) 2 in a desiccator at room temperature.

Mechanical properties

Standardized dumbbell-shaped specimens (ISO 527-2) with a 13 mm long and 3.3 mm wide rectangular working section were cut from the films with a cutting die. These were then tested for tensile resistance on a TA Instrument DMA with a force ramp of 1 N min K1 . This gave ultimate tensile strength (UTS), elastic modulus (E) and the elongation at break (EL); the film toughness was obtained from the area of the stress-strain curves, and the tensile strength-elastic modulus ratio UTS/E also obtained [29]. Five to seven measurements were made for each type of film. Dynamic mechanical analyses (DMA) of the films were performed over the temperature range K120 to 20 8C with a scanning rate of 3 8C min K1 . The static force, amplitude and frequency were maintained constant at 0.2 N, 2 mm and 1 Hz, respectively. To avoid water condensation, all dumbbell specimens were covered with a layer of silicone grease, which does not affect the dynamic mechanical properties of starch-based materials [START_REF] Gaudin | [END_REF].

Freeze fracture

In order to examine their internal structure, the films were freeze-fractured using the following procedure. A 1 mm thick cylinder of film was obtained by casting multiple layers of film-forming latex/starch mixtures between two copper cups. The cylinders were then frozen at K150 8C and under reduced pressure before being fractured with a blade. The freeze-fractured surfaces were shadowed with platinum and then with carbon. The replicas were then isolated from the samples by washing in DMSO and acetone before being placed on a copper grid. Finally they were observed on a Philips CM120 TEM.

Contact angle with water

The films were kept at 54% r.h. before being tested on a Contact Angle Meter GBX (Romans sur Ise `re, France). Three drops of water were placed successively on a piece of each film, and their left and right angles measured. The value reported is the average of six measurements.

Water absorption

Pieces of film equilibrated at 54% r.h. were placed in a sealed container above a saturated solution of K 2 SO 4 (97% r.h.) for 24 h. They were weighed at regular intervals; massgain values reported below are the average of three measurements.

Results and discussion

Mechanical properties

The properties arising from using unmodified NRL are first considered. As noted recently [30], adding NRL to starch-based materials without any external plasticizer does not result in any improvement of the material properties. It was found in the present work that it is, however, possible to add up to 20% of NRL without any apparent phase separation, this being presumably due to the natural stabilizers in NRL (proteins and phospholipids). With the further addition of 20% (w/w with respect to dry starch) of glycerol, one can add as much as 30% NRL without phase separation. This amount of glycerol acts as a plasticizer, and the formed films are supple enough to be easily handled, even without the addition of latex. Mechanical properties of the formed films are shown in Table 2. Addition of from 5 to 20% NRL caused a decrease of both tensile strength and elastic modulus but no modification of the elongation at break, of film toughness or of UTS/E (which quantifies the resistance of the films to crack [29]). Adding 30% NRL to the glycerol-plasticized starch films resulted in an improvement of the film plasticity: the tensile strain increased to 11% with no significant change in tensile strength or elastic modulus. The rubbery phase now appears to interact significantly with the starch matrix, rather than merely acting as an inert filler.

The effects of addition of modified (grafted) NRL are now examined. It is expected that the addition of modified natural rubber, with the hydrophilic hairs likely to form hydrogen bonds with the starch, would enhance compatibility between the hydrophobic component (polyisoprene), to which they are grafted, and the hydrophilic (starch) component. When considering the properties of the effects of addition of modified NRL, it is also necessary to consider which effects may arise from ungrafted polyDMAEMA alone.

The interpretation of the observations is made complex by the presence of both grafted and ungrafted poly-DMAEMA. Unfortunately, it proved impossible to separate the free chains from the latex, presumably because the free chains tend to strongly associate with grafted poly-DMAEMA on the particle surface. Attempted separation using centrifugation led on the one hand to an unrealistically low value of the fraction of unbound chains and moreover centrifugation resulted in irreversible coagulation of the latex. However, the total amount of free chains could never exceed 4.5% of the starch weight (for an addition of 30% of GRL15), and such a low ratio of external plasticizer is very unlikely to be responsible for the large change in properties observed.

This difficulty in principle could be overcome by adding free polyDMAEMA to NRL and carrying out the same tests. However, this leads to coagulation of the NRL if one adds polyDMAEMA (which is alkaline) and readjusts the pH to 7.

The addition of modified natural rubber with 5% DMAEMA did not result in any visible difference in the mechanical properties the formed film, compared to those for the same amount of unmodified NRL, if the pH was such that the polyDMAEMA chains were uncharged (GRL5(n)). However, when the pH of the latex was decreased to 7 (by addition of hydrochloric acid before being added to the starch/glycerol solution), the tensile properties of the films were significantly changed. With a GRL5 content of 15%, the tensile strength and the elastic modulus dropped to 6.3 and 8.5 MPa, respectively, and the strain at break reached 16.2%, while with addition of 30% of GRL5(n) the strain at break was only 12.2% (Table 2). This trend continued when the GRL5 content was increased, with an increase in the elongation up to 38.1% and a large decrease in both tensile strength and elastic modulus.

At pH 7, there will be a significant number of positively charged units on the polyDMAEMA chains, as the pK a of the monomer is w9.4 [31]. These charged groups, many of which are on chains covalently linked to the polyisoprene [16], should increase the interactions between starch and rubber and would be expected to lead to a significant change in the mechanical properties of the composite material. Indeed as shown in Table 2, the resulting films have a higher elasticity, high toughness (even for a small amount of latex), and high resistance to crack. Although the modified latexes contain significant amounts of ungrafted polyDMAEMA chains, the preceding argument suggests that the observed effects are dominated by grafted chains. It is highly unlikely that ungrafted chains could have a sufficient plasticizing effect to explain this significant change in material properties. The amount of polyDMAEMA compared to starch is low: e.g. a film containing 20% of GRL10 contains only 2% of polyDMAEMA. Literature data [32] show that much higher fractions are needed to affect significant property change in a blend (i.e. in the absence of grafting).

The observed effect of the amount of polyDMAEMA in the parent NRL shows that 10% gives good compatibility between the two phases. The addition of 20% of GRL10 gives high strain at break and the highest toughness (Table 2). With 30% of GRL10, the specimens were too rubbery to be tested: they exceeded the maximum elongation acceptable by the apparatus. NRL and GRL5(n) give relatively small decreases in the elastic modulus, while GRL5 and GRL10 lead to the greatest decrease, with GRL15 being intermediate (Table 2). This could be due to a larger amount of ungrafted water-phase polyDMAEMA in GRL15. The percentages refer to the weight of starch and latex. All films contain 20% of glycerol w/w starch mass.

Dynamic-mechanical properties

The dynamic mechanical properties of a polymer blend give information on the degree of phase mixing. If the blends are immiscible, the tan d curves will show the presence of two damping peaks corresponding to the T g s of the of the individual polymers [33]. If there is complete miscibility between the two polymer phases, only a single peak will be seen, lying between the T g s of the component polymers. Broadening and shifts of the damping peaks are characteristic of partially miscible systems [34]. In the present system, any blending between the two main phases, starch and polyisoprene, will only be that induced by the relatively small amount of polyDMAEMA, which clearly could not lead to the sort of intimate blending seen in some other modified NRL systems where there is extensive grafting between both main phases [35]. It is important to recall that the maximum in tan d in DMA at some particular frequency (the value of T d is of course frequencydependent), which gives the dynamic transition temperature T d , is less than the value of T g as measured by DSC. This is because each technique measures different properties. DSC measures the change in heat capacity going from the 'frozen' to the 'unfrozen' chain, whereas DMA measures the change in mechanical response of these chains [START_REF] Ferry | Viscoelastic properties of polymers[END_REF][START_REF] Wunderlich | Thermal Analysis[END_REF]; the two techniques give the effectively different averages of the complex dynamics of chain motion.

In the present system, there will be two T g s and hence two T d values in the parent polymers: those for (modified) natural rubber and the glycerol-plasticized starch. Fig. 1 shows the observed temperature dependence of tan d at 1 Hz between K100 and 0 8C for the reference starch/glycerol and a range of GRL blends with different amounts of polyDMAEMA, while Fig. 2 shows the corresponding loss moduli. The T d for the reference film is ca. K48 8C, and similar values are obtained with the addition of NRL and of GRL5(n). The peak corresponding to the T g of cis-polyisoprene is seen at K63 8C, and there is no significant change in the T d for the starch, indicating insignificant interaction between the phases. When the pH of the GRL is such that the hydrophilic polyDMAEMA chains are charged, the thermo-mechanical behavior of the formed film is very different. The starch T d increases to K38 and K 32 8C with the addition of 20% of GRL5 and GRL10, respectively. This can be ascribed to strong interactions between the starch and the polyDMAEMA chains, whose T g is w10 8C [START_REF] Martin-Gomis | [END_REF]. This interaction extends to the rubber phase, as the NR peak becomes a small shoulder and the starch peak is somewhat broadened.

Film morphology

The freeze fracture of thick specimens of films containing 30% of NRL and GRL10 gives information on their internal structure, although it must always be borne in mind that any freeze-fracture micrograph is easily over-interpreted. For unmodified NRL, the latex particles are clearly visible in the starch matrix. The adhesion between the two phases is imperfect: some cracks appear on some of the particles edges (panel A of Fig. 3). With the grafted latex GRL10, containing 10% of polyDMAEMA, the internal structure of the film is clearly different. The particles are completely embedded in the matrix (panel B of Fig. 3). This is consistent with good compatibility between the fillers and the matrix induced by the presence of grafted cationic hairs.

The particles in panel C are larger than those in panel B, which is because of an adventitious selection on the TEM grids: NRL has a very broad range of particle sizes (from w0.1 to 2 mm). It is also noted that adventitious selection of a grid area can lead to the erroneous impression that the GRL 10 blend micrographs indicate a higher amount of latex than for NRL. In fact the overall amounts are the same in all three samples. It is well known (e.g. [39]) that quantitative information from such micrographs requires examination of w10 3 particles.

Hydrophobicity

The addition of hydrophobic particles to a hydrophilic component should enhance the water resistance of the resulting film, but this may be countered by the presence of the hydrophilic polyDMAEMA. Results for contact angle are shown in Fig. 4, and for water adsorption in Fig. 5. The contact angles are greater than are typical for starch-based materials, which is not unsurprising given that a hydrophobic component is being added. With pure NRL the change is obvious. With up to 15% NRL, the contact angle with water increases for both the film face in contact with the air during the casting and for the other face, in contact with the polystyrene petri dish. With an addition of 20% of NRL, the angle tends to decrease, which possibly could be due some coagulation resulting in a weaker dispersion of the fillers in the matrix.

The addition of NRL renders the resulting film more hydrophobic, and also results in a net decrease of the water absorption of the material in a high-humidity atmosphere. After 24 h at 97% r.h. the mass gain of a film containing 20% of NRL is only 67% compared to the 88% of the reference film.

The chemical modification of NRL by DMAEMA causes the opposite effect even if the ratio of hydrophilic chains in GRL10 is only 10%. The contact angle decreases on both faces, and the water absorption greatly increases. This is because the cationic polyDMAEMA chains are even more hydrophilic than the starch and the interactions between them are not sufficiently strong to enhance the water resistance of the film.

Explanations for the dip in the curve of the water contact angle of the surface facing the polystyrene of the film containing 5% of GRL10 (which is reproduced with three separate samples) and water contact angle difference between the two faces of films containing modified latex have not yet been elucidated.

Conclusions

Films formed from the addition of two types of rubber latexes (natural rubber latex, and natural rubber latex grafted with 'hairs' of a cationic hydrophilic polymer, polyDMAEMA, the latter latex also containing ungrafted polyDMAEMA) to a glycerol/starch solution show a range of physical properties which can be explained by the competing effects of the hydrophilicity of the starch and the polyDMAEMA, the hydrophobicity of the NRL, and hydrogen bonding between the grafted polyDMAEMA and the starch. NRL acts essentially as inert filler: the tensile strength and modulus decrease and the elongation slightly increases. The adhesion between the phases is imperfect, and there is an increase in the hydrophobicity of the starch film.

The DMAEMA-modified NRL significantly changes the properties of the starch/latex films. In the cationic form, the polyDMAEMA chains enhance greatly the compatibility between the rubber and starch phases, with an optimum compatibility obtained for about 10% (w/w) DMAEMA relative to rubber. This leads to a large increase in film elongation and toughness, but also in an increase of the water sensitivity of the composite material.

These results have potential applications in improvement of material properties of materials made from judicious addition of modified and unmodified NRL to polysaccharide-based agricultural 'waste' material.

Fig. 1 .

 1 Fig. 1. Temperature dependence of damping ðtan dÞ for films containing 20% of the different latexes. The reference film comprises only starch and glycerol.

Fig. 2 .

 2 Fig. 2. Temperature dependence of loss moduli for films containing 20% of the different latexes. The reference film comprises only starch and glycerol.

Fig. 3 .

 3 Fig. 3. TEM micrographs of freeze-fractured surfaces of films containing NRL (A) and GRL10 (B and C). Latex content is 30%. The arrows on panel A show some of the cracks on particles edges.

Fig. 4 .

 4 Fig. 4. Average contact angle and standard deviations for water on both surfaces of films, as a function of latex content.

Fig. 5 .

 5 Fig.5. Water absorption at 97% r.h. of films without latex, with 20% of NRL and with 20% of GRL10.

Table 1

 1 Recipes for grafted natural rubber latexes (GRL), final solid content, and amount of DMAEMA relative to the total mass of solids

		GRL5	GRL10	GRL15
	NRL/g	51.85	52.07	51.96
	DMAEMA/g	1.604	3.208	5.238
	CHP/g	0.069	0.139	0.205
	NH 3 (2.5%)/g	100.7	100.1	100.1
	TEPA/g a	0.095	0.179	0.270
	Solid content	21.1	22.1	23.2
	(%)
	DMAEMA	3.0	5.8	9.1
	content (%)

a Made up to 10% solution with distilled water.

Table 2

 2 Average value and standard deviation (in parentheses) of the ultimate tensile strength (UTS), strain at break (EL), elastic modulus (E), toughness and UTS/E value of the different films conditioned at 54% relative humidity

		Latex content (%	Ultimate tensile	Strain at break (%)	Elastic modulus E	Toughness	UTS/E
		w/w)	strength UTS		(MPa)	(MJ m K3 )
			(MPa)
	Reference	0	12.3 (1.2)	4.0 (0.4)	46.0 (6.9)	0.33 (0.05)	0.27 (0.07)
	NRL	5	10.3 (0.4)	4.1 (0.7)	33.1 (1.9)	0.26 (0.03)	0.31 (0.03)
		10	10.1 (0.8)	3.5 (0.7)	30.8 (3.5)	0.25 (0.02)	0.33 (0.06)
		15	8.4 (0.7)	4.4 (0.6)	27.1 (2.2)	0.26 (0.01)	0.31 (0.05)
		20	7.9 (0.9)	4.6 (0.4)	23.6 (3.6)	0.27 (0.01)	0.33 (0.11)
		30	8.0 (0.7)	11.0 (1.0)	19.5 (2.1)	0.7 (0.1)	0.41 (0.1)
	GRL5(n)	5	9.8 (1.2)	4.2 (0.3)	31.6 (5.2)	0.28 (0.01)	0.31 (0.12)
		10	9.6 (1.4)	3.4 (1.1)	36.7 (5.9)	0.31 (0.06)	0.26 (0.1)
		15	8.6 (0.9)	5.9 (1.2)	21.5 (2.1)	0.32 (0.03)	0.4 (0.08)
		20	8.9 (0.9)	5.3 (0.7)	23.2 (0.8)	0.30 (0.06)	0.38 (0.06)
		30	6.4 (0.7)	12.2 (2.2)	15.6 (3.6)	0.58 (0.06)	0.41 (0.18)
	GRL5	5	11.0 (0.3)	5.4 (0.4)	23.6 (2.4)	0.3 (0.1)	0.47 (0.06)
		10	9.6 (2.0)	6.8 (2.7)	19.6 (2.4)	0.3 (0.1)	0.49 (0.18)
		15	6.3 (0.9)	16.2 (6.2)	8.5 (2.5)	0.6 (0.2)	0.74 (0.46)
		20	4.8 (0.4)	20.4 (3.0)	5.8 (0.2)	0.6 (0.2)	0.83 (0.1)
		30	2.2 (0.1)	38.1 (0.9)	1.6 (0.5)	0.59 (0.06)	1.37 (0.72)
	GRL10	5	9.8 (0.2)	4.3 (0.1)	25.7 (3.1)	0.33 (0.05)	0.38 (0.06)
		10	9.1 (0.9)	9.3 (4.6)	20.2 (4.9)	0.5 (0.2)	0.45 (0.2)
		15	7.1 (0.3)	14.2 (0.7)	13.1 (0.4)	0.75 (0.06)	0.54 (0.04)
		20	3.1 (0.1)	39.9 (8.3)	2.8 (0.7)	0.9 (0.1)	1.11 (0.41)
		3 0	-	-	-	-	-
	GRL15	5	8.2 (0.6)	3.3 (1.4)	28.6 (3.8)	0.13 (0.03)	0.29 (0.06)

Acknowledgements AR gratefully acknowledges the support of a Linkage-International Fellowship from the Australian Research Council. We also appreciate the provision of facilities and expertise by the University of Sydney Electron Microscope Unit, and of the assistance of Rachelle Ward in measuring the amylose content of the rice starch. The Key Centre for Polymer Colloids is established and supported under the Australian Research Council's Research Centres Program.


Sehmi Saad
email: sehmi.saad@gmail.com

Fayrouz Haddad
email: fayrouz.haddad@im2np.fr

Aymen Ben Hammadi
email: benhammadiaymen@gmail.com

A 270 Hz, Fine Frequency Tuning Class-C Oscillator Using Capacitive-Inductive Degeneration Technique in 130 -nm CMOS

Keywords: Capacitive-Inductive Degeneration Coupling, Class-C oscillator, LC-VCO, robust start-up, CMOS I

This paper presents a high-resolution Class-C type voltage-controlled oscillator (VCO) with a robust startup. The proposed oscillator achieves a minimum frequency quantification step of 270 Hz without adding any additional dithering approach. The very fine tuning is obtained through a capacitive-inductive degeneration coupling (CIDC) technique which reduces the capacitance value without appreciably affecting the intrinsic oscillator phase noise (PN). This technique also provides a negative transconductance to compensate for losses in the LC tank, resulting in high oscillation amplitude and relatively low power consumption in the oscillator core. The proposed design, implemented in 130 nm CMOS technology, achieves 45% tuning range. It exhibits a phase noise of -107 dBc/Hz @ 1-MHz offset, while it draws 5.75 mA from a supply of 1.2 V, resulting in a FoM of 165 dBc/Hz.

INTRODUCTION

Nowadays, combining low power consumption, high speed and high spectral purity in voltage-controlled oscillator (VCO) is one of key challenges in CMOS circuit design and process methodologies of radio frequency (RF) VCO. Although the progress in technologies evolves by itself the enhancement of these goals, the target resolution for modern wireless applications is still quite challenging to reach fine frequency steps.

Several techniques such as the capacitive divider [START_REF] Chen | A 9 GHz dual-mode digitally controlled oscillator for GSM/UMTS transceivers in 65 nm CMOS[END_REF], the capacitive dithering [START_REF] Staszewski | A digitally controlled oscillator in a 90 nm digital CMOS process for mobile phones[END_REF] and the capacitive degeneration [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF][START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF][START_REF] Saad | A fine frequency tuning for LC-DCO using enhanced capacitive degeneration technique[END_REF] have been suggested to minimize the unit step capacitance in LC oscillators. For example, a capacitive divider network that can be switched in parallel to the LCtank is used in [START_REF] Chen | A 9 GHz dual-mode digitally controlled oscillator for GSM/UMTS transceivers in 65 nm CMOS[END_REF] to obtain a reduction of the minimum effective capacitance. This method can improve the oscillator frequency resolution but the sensitivity to parasitics and mismatches limits the robustness of the final design. A more reliable idea has been developed by Staszewski et al. in [START_REF] Staszewski | A digitally controlled oscillator in a 90 nm digital CMOS process for mobile phones[END_REF]. It consists in dithering least significant bits (LSB) of the oscillator control word (like ΣΔ modulator). The technique can achieve fine frequency from 12 kHz to 30 Hz, but as it occurs in any ΣΔ data converter, this solution faces problem of the quantification noise moved to higher frequencies. Another attractive approach is proposed by Fanori et al. [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF]. It consists in moving a part of the capacitor bank from the LC-tank to the sources of the cross-coupled transistors of the oscillator. equivalent scheme [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF] At this level, an intrinsic shrink factor reduces the equivalent capacitive element and lowers considerably the equivalent oscillator frequency resolution (an averaged resolution of 300 Hz in [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF], with a reasonable noise performance at the cost of several design constraints). The topology requires both small transconductance and high capacitance value to ensure an adequate shrink factor [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF] and, at the same time, a large transconductance to sustain the VCO oscillations. That increases the design complexity, power consumption and area occupation.

In [START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF], we proposed an enhancement to overcome the limitations of [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF]. By placing a resistive element in parallel to the degenerated capacitance, a robust start-up even is ensured with a relatively large choice of capacitance values, resulting in 3 kHz frequency resolution. In this paper, an RLC stage through dithering process is designed and implemented in a LC oscillator. Alike to [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF] and [START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF], the topology exploits small degenerated capacitance with more enhanced shrink effect, able to achieve a fine frequency of some hundred hertz without any additional circuit.

The paper is organized as follow. Section II introduces the basic idea of capacitive degeneration and its enhancement with a resistive element. Section III describes the amended topology with a detailed analysis. To verify the theoretical affirmations, an overview of implementation in a Class-C oscillator is proposed and followed by the simulated results in Section IV. Conclusions are given in Section V.

II. CAPACITIVE DEGENERATION TECHNIQUE

A. Basic idea using only a capacitor

The basic idea of minimizing the effective capacitance is innovated by Fanori et al. [1] [6]. As shown in Fig. 1(a), the technique consists in moving a part of the varactor bank from the LC-tank to the sources of two transconductances that built the negative impedance of the oscillator. The result shows that the moved portion of capacitor is submitted to an effort of shrinking factor effect appearing in the capacitive degeneration of transistors M1 and M2. Since the circuit works in large signal regime, the frequency tuning characteristics can be evaluated, as well as the shrink factor through small signal time variant analysis [START_REF] Mazzanti | Class-C Harmonic CMOS VCOs, With a General Result on Phase Noise[END_REF]. In order to simplify the analysis, it is supposed that the transconductance of cross-coupled transistors, gm, is averaged over a time interval equal to one period of the oscillation frequency. Under the assumptions, the equivalent scheme reported in Fig. 1(b) can model the oscillator reported in Fig. 1(a).

The effect of the moved capacitance Cdgn on the oscillator frequency tuning characteristics can be determined by evaluating the admittance Y as m eq eq Y g jC   :

22 2 2 2 2 2 2 2 4 2 4 4 dgn LO m m LO dgn m dgn LO m dgn LO C g g Y j C g C g C            ( 1
)
where ωLO is the proper oscillation frequency of the oscillation and Cdgn is negative since it follows the sign of current that crosses M1 and M2 (due to the drain-gate connection).

For Cdgn >> gm/2ωLO, Eq. ( 1) can be simplified as follows:

2 2 2 2 2 m m m LO dgn LO dgn f dgn LO g g g Y j C j C Q C                  ( 2
)
where Qf is the quality factor of the impedance which models the degenerated effect in the switch defined by:

    2 f m osc dgn Q Z Z g C      (3)
gmeq becomes the classical negative transconductance, which restores tank losses [START_REF] Rohde | The Design of Modern Microwave Oscillators for Wireless Applications: Theory and Optimization[END_REF], while Ceq is the capacitor Cdgn shrunken by a factor Fshrk equal to

  2 2 2 shrk f m dgn LO F Q g C     .
Starting from Eq. ( 2), gmeq and Ceq have been plotted in Fig. 2 versus Cdgn. The equivalent capacitance Ceq shows a nonmonotonic behavior, where it starts to decrease for Cdgn>gm/(2ωLO), whereas the equivalent conductance gmeq shows a monotonic behavior that asymptotically tend to -gm/2. As well, the plot in Fig. 2 indicates that in order to sustain oscillation, the absolute value of the real part of Y, should be larger enough than 1/Rloss, but also sufficiently small to keep the required effect of reducing the effective capacitance. This trade-off does not allow a degree of freedom, since Cdgn and gm are univocally defined from Eq.

(3), i.e. if the Q-factor of the resonator becomes low, the transconductance gm becomes large leading to excessively large value for capacitor Cdgn. Contrarily, if Cdgn<gm/(2ωLO), the real part of Y tends dramatically to zero and losses cannot be more compensated. To resolve this drawback, an enhanced capacitive degeneration using resistive element is presented in the next sub-section.

B. Enhanced capacitive degeneration using a resistive element

As mentioned previously, placing in the region where Cdgn >>gm/2ωLO, cannot provide a required shrink factor owing to weakness value of gmeq. Fig 3 shows a possible enhancement to compensate losses in the tank. A resistor Rdgn is added in parallel to the degenerated capacitor bank. Thus, the shrinking effect will be controlled accordingly to the following equation:

m m m R R C R g g Y jC R R C R R C g g g g                                            (4)
dgn dgn LO dgn m m R C R g g                and   2 2 2 dgn dgn LO m m R C g g  
and dgn R ,the expression of the admittance Y becomes:

1 2 2 dgn dgn LO dgn dgn m m R Y jC R R g g             (5)
Comparing this expression with Eq. ( 2), the main difference is that in this structure, gmeq becomes the conventional negative impedance placed in series with the resistance Rdgn, while Ceq is equivalent to Cdgn shrunked by a factor Fshrk given by:

2 2 2 2 2 2 2 2 dgn dgn shrk dgn dgn dgn LO dgn m m m R R F R R C R g g g                         (6)
This factor is equal to (Qf/CdgnRdgnω)², where Qf is the Qfactor of the admittance which models the degenerated effect in the feedback, and shown to be:

2 [] 2 [] dgn dgn LO f dgn m CR Y Q Y R g      (7)
Similar to [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF], to decrease the equivalent capacitance, Qf could be further decreased to minimize the frequency resolution of the oscillator. Based on Eq. ( 5), the transfer function of gmeq and Ceq are plotted in Fig. 4, for same parameters as in [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF]. It can be seen that a different variation of gmeq over Cdgn is obtained. The curve starts from 1/gm+Rdgn and has almost a linear behaviour asymptotically close to -4.5 mS. The advantage of using an additional resistance is that the real part of Y is considerably retained for low value of Cdgn, because Rdgn will play the same role of gm, and is converted to transconductance. This results in sustaining the oscillation although small values of Cdgn are used. Another advantage is that Ceq is optimized and shows quite linear tendency which is beneficial to improve frequency resolution and tuning range.

The minimum resistance Rdgn(min) must be chosen to satisfy the required shrink factor. It is determined by the following formula:

2 2 1 2 2 dgn shrk dgn dgn dgn LO m m R F R R C g g                  (8)

III. PROPOSED CAPACITIVE-INDUCTIVE DEGENERATION COUPLING TECHNIQUE

In the proposed circuit presented in Fig. 5, a couple of resistor and inductor is placed in parallel to the variable capacitor bank.

To study the proprieties effect of changing the capacitance C and the insertion of R-L components on the tuning of the oscillator frequency characteristic, the same methodology mentioned earlier has been applied. The admittance Y in the equivalent scheme of Fig. 6 is given by:

L R R L C R L R L L R g Y jC R L R L R L C R L R L R L g g g g                                          2 2 n dgn dgn C        (9) Since 2 2 2 dgn dgn dgn LO mN R L C g         2 2 ( ) dgn dgn LO dgn dgn LO mN R L R L g          and   2 2 2 4 dgn dgn dgn LO R L C     2 dgn LO dgn L R   ,
Eq. ( 9) can be rewritten as follow:

  2 22 2 2 3 2 2 2 ( ) 2
L R R L R L LR g Y jC R L R L R L R L g g                               (10)
Referring to the equation mentioned in Fig. 6, the first term is the negative conductance given by:

  2 2 2 2 2 ( ) 2
L R R L R L g g R L R L g                  (11)
The second term reveals the equivalent capacitance Ceq emulated in the capacitance Cdgn shrunked by a factor of:

LR F R L R L R L C g g                     2 2 2 2 2 () dgn dgn LO dgn dgn LO dgn dgn LO m LR R L R L g        (12)
Further, the quality factor Qf of Y where its models the degenerated effect in the feedback, can be expressed as: The result shows that gmeq becomes more steady throughout capacitance Ceq. Practically, it has a horizontal straight behavior of around -5 mS. Considering the behavior of Ceq, it has a wide linear monotonic-curve without much variation on the value of shrunken capacitor. The shrink factor is proved in Fig. 8, with results from [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF] and [START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF] for the same conditions. From the different plots of Fig. 8, it can be seen that a capacitor of 2 pF is reflected in the parallel to the LCtank into an equivalent capacitance of 1.62 fF with a shrinking factor 1/Fshrk of about 1233. This factor is, respectively, 14 and10, times greater than factor in [START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF] and [START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF]. Prop. work Ref.

  2 2 3 2 [] [] 2 dgn dgn dgn LO f dgn dgn LO dgn dgn LO mN C L R Y Q Y R L R L g             (
[4] Ref.

[3] 1233 For Cap.

= 2pF

Fig. 8. Shrink factor versus capacitance Expression of the resonance frequency ωLO of the proposed circuit (cf. Fig. 6), as a function of the capacitance Cdgn placed between the sources of the negative resistance of NMOS transistors, can be found and deduced by the following expression:

0 ( ) 1 2 dgn LO dgn f TANK C f C Q C       ( 14
)
where Qf is given by Eq. ( 13) and

TANK TANK L C ω 1 0 
represents the resonant frequency of the conventional LC oscillator.

Eventually, the shrink factor is influenced by Ldgn, Rdgn and ωLO. Eq. ( 15) can be settled straightforwardly to an inequality < 1 in order to keep the effect at optimum:

2 2 2 2 mN LRω F 1 2 ( ) g dgn dgn LO shrk dgn dgn LO dgn dgn LO R L R L            (15)
IV. APPLICATION TO A CLASS-C OSCILLATOR In order to validate the presented equations, the VCO depicted in Fig. 9 is designed in 130 nm CMOS technology with an NMOS varactor and a 32-bits fine array. The proposed architecture consists of a complementary Class-C oscillator as the core, a bias circuitry, and coupling stages with capacitive source degeneration. The proposed circuit was tailored to RF ISM frequency band applications, with a center frequency of XX GHz.

The tank consists of a spiral differential inductor of 3.75 nH with a simulated maximum quality factor of 34.6. For the coarse tuning bank, a NMOS varactor providing an average gain of around 1.025 GHz/V was implemented. The fine tuning bank was realized as reported in Fig 9 . Since it is difficult to approve the effect of unit capacitance of a few atto-farad on the oscillation frequency, 32 capacitor pairs are tied in parallel. The row of varactor pairs is implemented with pair NMOS transistors in the inversion mode configuration. Although the C-V characteristic is non-linear and near monotonic, the transition from Cmim to Cmax is very biased which is effectively useless tuning gain in analog domain, so a digital toggling command is adopted to simplify the control of the row. The untied unit varactor pair is controlled by the Tuning Control Word (TCW) where it measures a unit per one digital weight variation capacitance of 258 aF while changing the control voltage from 0 to 1 V. The unitary NMOS elements size is W=10 µm and L=0.13 µm with a maximal total capacitance of 8.3 pF.

To create the negative gm, a cross-coupled pair is adopted, and the complementary topology (NMOS and PMOS) is used to reduce power consumption. To operate in Class-C, transistors of the cross-coupled pair are both biased in the saturation region. Hence, the bias voltage of NMOS (Vgsn) has to be lowered with respect to the supply voltage. Although this choice is not mandatory, it increases the transconductance of the differential pair and at the same time reduces the phase noise [START_REF] Hong | Robustness Enhancement of a Class-C Quadrature Oscillator Using Capacitive Source Degeneration Coupling[END_REF]. In addition, a large capacitance (Ctail) is connected at the drain of the tail current source, which filters the high-frequency noise injected by the bias and aids in shaping the injection current for a given shrinking factor [START_REF] Fanori | Highly Efficient Class-C CMOS VCOs, Including a Comparison with Class-B VCOs[END_REF].

The simulated fine-tuning frequency characteristic, shown in Fig. 10, demonstrates a very good agreement with the theoretical analysis. It has the same linear behavior as that of the theoretical curve. In this slope of operation, the shrink factor is around 1233 (0.82%) with a fine-tuning range of about 8.652 kHz, under 32 varactors. The attained effective fine-frequency step is approximately 270 Hz through capacitance-inductive degeneration.

In Fig. 11

P mW f FoM L f f mW              (16
) where L(Δf) is the measured phase noise in dBc/Hz at a frequency offset of Δf from the carrier frequency f0 and PDC is the measured power dissipation of the oscillator. In comparison with performances of other works, as shown in Table I, the proposed oscillator achieves the superior frequency resolution and the lowest power dissipation at 3.2 GHz with reasonable phase noise even in a 130 nm CMOS technology. Capacitive degeneration, (b) Without dithering, (c) Tunable acting on current variation V. CONCLUSION A capacitive-inductive degeneration coupling technique (CIDC) has been proposed and applied to a CMOS Class-C oscillator. By using an RLC circuit mounted in parallel to the tank, capacitance is shrunken and frequency step can be further reduced. By exploiting a PMOS varactor in inversion region C-V characteristics, the equivalent capacitance diminishes to approximately 0.21 aF. This allows to perform a minimum frequency step with a resolution of about 270 Hz without any dithering process. The small signal analysis demonstrates the limits and the potentialities of the proposed idea, while the simulation results prove its robustness in 2.2-3.2 GHz frequency band. Even if the proposed structure is implemented in 130 nm CMOS technology, the oscillator has superior performances in terms of frequency tuning resolution with reasonable power consumption and phase noise compared to the stateof-the-art.

Fig. 1 .

 1 Fig. 1. (a) Simplified bloc diagram of the oscillator, (b) oscillatorequivalent scheme[START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF]

Fig. 2 .

 2 Fig. 2. Real and Imaginary parts of admittance Y (gm=10 mS, Rloss=250Ω, LO=3.6 GHz)[START_REF] Fanori | Capacitive Degeneration in LC-Tank Oscillator for DCO Fine-Frequency Tuning[END_REF]

Fig. 3 .

 3 Fig. 3. (a) Simplified bloc diagram of the oscillator, (b) oscillatorequivalent scheme[START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF]

Fig. 4 .Fig. 5 .

 45 Fig.[START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF]. Real and imaginary parts of admittance Y vs capacitance (gm=10 mS, Rdgn=455 Ω, Ldgn=025 pH, f0=3.6 GHz)[START_REF] Saad | A 5-Mw, 1.2-3.5-Ghz capacitive degeneration in LCdigitally-controlled oscillator for nano-satellite frequency synthesizers in 90-Nm CMOS[END_REF]

  13) To confirm the effect of adding Rdgn and Ldgn, gmeq and Ceq are theoretically extracted in Fig 7, while maintaining the same configuration for fair comparison.

Fig. 7 .

 7 Fig. 7. Real and imaginary parts of admittance Y vs capacitance (gmN=10 mS, Rdgn=455 Ω, Ldgn=0.25 pF, f0=3.6 GHz) of the proposed structure

Fig. 12

 12 Fig. 12 depicts the simulated phase-noise characteristic from 100 Hz to 10 MHz of offset frequency. The oscillator achieves a phase noise of -107 and 120 dBc/Hz at 1 and 3 MHz offset, respectively, from the carrier frequency when oscillation frequency is 2.23 GHz. The Figure of Merit (FoM) evaluated by Eq. (16) according to [11] is equal to 166 dBc/Hz.     0 10 10 -20log 10 log 1 DC

Fig. 11 .Fig. 12 .

 1112 Fig. 11. Oscillator INL/DNL vs input code and 1 LSB for fine-tuning bank

  , the obtained Integral-Non-Linearity (INL) and Integral-Non-Linearity (DNL) curves are reported. A good linearity is achieved since the DNL is 0.37 LSB while the INL is 0.42 LSB.

								5	µm
								0.13 µm	Source follower
	Zin		Ifb
		Mfb	M1					M2
	Vctr2
	M3							L
							2C	Vctr-g	2C
			Vout+						Vout-
								Cvar
				M3				Vgsn	M4
								Ldgn
								FCW
	VCO Core
								Cdgn
	10	µm
	0.13 µm		Itail				Rdgn	Itail
			Ctail	M5			Vbias		Vbias	M6	Ctail
		Fig. 9. Schematic of the proposed Class-C type degenerated VCO
			2.233032
				2.2330305
		[GHz] Frequency	2.233024 2.233026 2.233028 2.233030	2.2330300 2.2330280 2.2330285 2.2330290 2.2330295	20	24	28	32	Simulation Theoretical Ideal line
			2.233022
				0	4		8 12 16 20 24 28 32
								TCW Code

Zoom Fig.

10

. Oscillator fine-tuning frequency response vs. the tuning control word

TABLE I .

 I OSCILLATOR PERFORMANCE COMPARISON

		This
	Parameter		[2]	[3]	[4]
		work
	CMOS Process [nm]	130	90	65	90
	Supply voltage [V]	1.2	1.4	1.8	1.0
	Tuning range [GHz]	2.23-3.25	3.2-4	2.9-3.6	1.22-3.52
	Frequency tuning
		45.7	22.2	26	65
	range (FTR) [%]
	Phase noise@1MHz
		-107	-136	-127.5	-105.8
	[dBc/Hz]
	Structure Cap.Tech.	Class-C	Class-B	Class-C	Class-B
				0.15-1.5 (a)
	Resolution [kHz]	0.27	12 (b)		3-16 (a)
				(c)
	Power [mW]	6.9	18	16	4.8-5.2
	FoM [dBc/Hz]	-165.6	N/A	-183	-174.8
	(a)


Tuncay Koç

Florencia Marchini

Gwenaëlle Rousse

Romain Dugas

Jean-Marie Tarascon
email: jean-marie.tarascon@college-de-france.fr

In the Search for the Best Solid Electrolyte-Layered Oxide Pairing for Assembling Practical All-Solid-State Batteries

Keywords: Solid-state battery, solid electrolyte, sulfides, halides, solvent-based synthesis, solvent-free synthesis, coated layered oxide, cathode composite

All-solid-state batteries (ASSBs) are the subject of large enthusiasm as they could boost by 50% the energy density of today's Li-ion batteries provided that several fundamental/practical roadblocks can be solved.

Focusing on the interface between the solid electrolyte and cathode active material (CAM), we herein studied the chemical/electrochemical compatibility of coated-layered oxide LiNi0.6Mn0.2Co0.2O2 with the three main inorganic electrolytes contenders, these being lithium thiophosphate (β-Li3PS4), argyrodite (Li6PS5Cl), and halide (Li3InCl6). Such electrolytes were prepared by either solvent-free or solution chemistry and paired with the CAM to form a composite which was further tested by assembling solid-state batteries using Li0.5In as negative electrode. Amongst the electrolytes prepared by dry routes, the best performing was found to be Li6PS5Cl followed by β-Li3PS4 and lastly Li3InCl6. In contrast, no general trend of benefit or detriment was observed when switching to solution route as it resulted in either performance improvement or deterioration depending on the electrolyte. Additionally, we show a strong dependence of the battery performance upon the presence of carbon additives. Lastly, we unraveled a pronounced chemical/electrochemical incompatibility of Li3InCl6 towards Li6PS5Cl and β-Li3PS4, hence questioning the design of hetero-structural cell architectures. Altogether, we hope these findings to provide guidance in the proper pairing of electrode-electrolytes components for designing highly performing solid-state batteries.

Introduction

All-solid-state batteries (ASSBs) are generating worldwide excitement among research institutions, battery makers and users as evidenced by the increasing volume of related scientific publications and daily announcements by various start-ups. This excitement was initiated back to 2011 by the discovery of an inorganic sulfide-based electrolyte (Li10GeP2S12) whose ionic conductivity of 10 -2 S/cm -1 was comparable to the best Li-based liquid electrolytes. 1 ASSBs are solvent-free devices, which offer safety advantages over lithium-ion systems while being theoretically capable of achieving a nearly 70% and 40% increase in volumetric (Wh/l) and gravimetric (Wh/kg) energy densities, respectively, taking advantage of Li-metal and bipolar cell configurations. [START_REF] Janek | A Solid Future for Battery Development[END_REF] However, for this to happen several roadblocks dealing with problematic electrode-electrolytes chemicalelectrochemical compatibilities at both negative and positive electrodes need to be equally addressed. [START_REF] Banerjee | Interfaces and Interphases in All-Solid-State Batteries with Inorganic Solid Electrolytes[END_REF][START_REF] Marchini | Li-Rich Layered Sulfide as Cathode Active Materials in All-Solid-State Li-Metal Batteries[END_REF][START_REF] Famprikis | Fundamentals of Inorganic Solid-State Electrolytes for Batteries[END_REF] Various strategies relying on Li-surface engineering or on the use of Li-alloying coatings (LixAu, etc.) acting as buffer interfaces have been tried. [START_REF] Yan | Selective Deposition and Stable Encapsulation of Lithium through Heterogeneous Seeded Growth[END_REF] The best results achieved so far have been obtained with "anodeless" ASSBs in which the Li-metal anode is replaced by a silver-carbon composite electrode. [START_REF] Lee | High-Energy Long-Cycling All-Solid-State Lithium Metal Batteries Enabled by Silver-Carbon Composite Anodes[END_REF] Positive electrodes of today's ASSBs are composites of a cathode active material (CAM, typically coated layered oxides) mixed with solid-ionic conductors and eventually carbon additives. [START_REF] Zhang | Interfacial Processes and Influence of Composite Cathode Microstructure Controlling the Performance of All-Solid-State Lithium Batteries[END_REF][START_REF] Strauss | Influence of Electronically Conductive Additives on the Cycling Performance of Argyrodite-Based All-Solid-State Batteries[END_REF] Their long cycling stability is presently plagued by chemical/electrochemical incompatibility issues between the oxide and the solid Li + conductor resulting in occasional decomposition reactions taking place at the interfaces. [START_REF] Walther | Visualization of the Interfacial Decomposition of Composite Cathodes in Argyrodite-Based All-Solid-State Batteries Using Time-of-Flight Secondary-Ion Mass Spectrometry[END_REF][START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] The degree of advancement of such decomposition reactions depends upon the quality of oxide coating but more so on the nature of the solid ionic conductor used. [START_REF] Xiao | Computational Screening of Cathode Coatings for Solid-State Batteries[END_REF][START_REF] Culver | On the Functionality of Coatings for Cathode Active Materials in Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Strauss | 2 ZrO 3 -Coated NCM622 for Application in Inorganic Solid-State Batteries: Role of Surface Carbonates in the Cycling Performance[END_REF] Nowadays, a wide panel of fast Li + conductors for solid-state battery applications are known. The lithium thiophosphate family is probably the largest and one of the most attractive for practical applications. It enlists typical structures such as β-Li3PS4 from the Li2S-P2S5 binary system, the argyrodite-type phases (Li6PS5X, X=Cl, Br and I) and the Li10GeP2S12 phases together with its derivatives obtained by partial substitution of Ge or S. Despite the latter exhibits the highest ionic conductivities, it is also the most unstable thermodynamic-wise and ruled out for practical ASSBs. [START_REF] Kato | Li 10 GeP 2 S 12 -Type Superionic Conductors: Synthesis, Structure, and Ionic Transportation[END_REF] More recently, halide-type ionic conductors Li3MX6 (M=In and trivalent rare earth metals Er, Y, etc., X=Cl, Br and I) have been revisited owing to their high ionic conductivity (10 -4 -10 -3 S/cm) [START_REF] Li | Progress and Perspectives on Halide Lithium Conductors for All-Solid-State Lithium Batteries[END_REF] and high stability against oxidation as deduced by density functional theory (DFT) calculations [START_REF] Park | Theoretical Design of Lithium Chloride Superionic Conductors for All-Solid-State High-Voltage Lithium-Ion Batteries[END_REF] and confirmed experimentally [START_REF] Li | Water-Mediated Synthesis of a Superionic Halide Solid Electrolyte[END_REF][START_REF] Zhang | Advanced High-Voltage All-Solid-State Li-Ion Batteries Enabled by a Dual-Halogen Solid Electrolyte[END_REF] . New electrolyte compositions such as Li2Sc2/3Cl4 and Li3ScCl6 have been recently uncovered. [START_REF] Zhou | A New Halospinel Superionic Conductor for High-Voltage All Solid State Lithium Batteries[END_REF][START_REF] Liang | Site-Occupation-Tuned Superionic Li x ScCl 3+ x Halide Solid Electrolytes for All-Solid-State Batteries[END_REF] Such richness of ionic conductors simultaneously represents a benefit and a hassle when selecting the best suitable solid electrolyte for pairing with a layered oxide positive electrode. Different electrolyte chemistries indeed differ in their thermodynamic stability against oxidation and reduction, in their electronic and ionic conductivity, in their softness (which is higher for halides than for argyrodite and thiophosphates) and lastly in their thermal stability. [START_REF] Kim | Material Design Strategy for Halide Solid Electrolytes Li 3 MX 6 (X = Cl, Br, and I) for All-Solid-State High-Voltage Li-Ion Batteries[END_REF] Moreover, their chemical-physical properties were also shown to depend on their synthesis method, either via solvent assisted reactions involving precipitation/crystallization in organic solvents, high temperature solid-state reactions (>500 °C) or mechanical ball milling followed by annealing steps. [START_REF] Kudu | A Review of Structural Properties and Synthesis Methods of Solid Electrolyte Materials in the Li2S -P2S5 Binary System[END_REF][START_REF] Ghidiu | Solution-Based Synthesis of Lithium Thiophosphate Superionic Conductors for Solid-State Batteries: A Chemistry Perspective[END_REF][START_REF] Schlem | Energy Storage Materials for Solid-State Batteries: Design by Mechanochemistry[END_REF][START_REF] Zhou | Solvent-Engineered Design of Argyrodite Li 6 PS 5 X (X = Cl, Br, I) Solid Electrolytes with High Ionic Conductivity[END_REF] Among them, the solvent-based routes appear to be particularly attractive for solid electrolyte (SE) -CAM composite preparation in a core-shell like structure (namely SE-coated oxide) by maximizing the energy density at electrode level while enhancing ionic percolation and boosting the cell performance. [START_REF] Wang | Interface-Assisted in-Situ Growth of Halide Electrolytes Eliminating Interfacial Challenges of All-Inorganic Solid-State Batteries[END_REF][START_REF] Yubuchi | Lithium-Ion-Conducting Argyrodite-Type Li 6 PS 5 X (X = Cl, Br, I) Solid Electrolytes Prepared by a Liquid-Phase Technique Using Ethanol as a Solvent[END_REF][START_REF] Park | Solution-Processable Glass LiI-Li 4 SnS 4 Superionic Conductors for All-Solid-State Li-Ion Batteries[END_REF] However, in practice, such solution processes were shown to lead to phase contaminations or surficial remaining solvent species on solid ionic conductors. [START_REF] Marchini | The Hidden Side of Nanoporous Β-Li 3 PS 4 Solid Electrolyte[END_REF][START_REF] Ruhl | Impact of Solvent Treatment of the Superionic Argyrodite Li 6 PS 5 Cl on Solid-State Battery Performance[END_REF] Most of these aspects have already been addressed for most of ionic conductors independently and almost no cross-comparison between such studies have been made, with the exception of a recent study on halides that report their stability against argyrodite. [START_REF] Han | Single-or Poly-Crystalline Ni-Rich Layered Cathode, Sulfide or Halide Solid Electrolyte: Which Will Be the Winners for All-Solid-State Batteries?[END_REF] Thus to fill this gap, we decided to embark ourselves in a benchmarking study of solid-state batteries using coated-LiNi0.6Mn0.2Co0.2O2 as part of a composite positive electrodes also including β-Li3PS4, Li6PS5Cl and Li3InCl6 as solid electrolytes. The latter were prepared either by solution or by solvent-free synthesis. For each electrolyte, we firstly studied the interplay between synthesis strategy, ionic/electronic conductivity and structural properties and later tested their performances in full all-solid-state batteries. The use of carbon additives together with the design of hetero-structural cell architecture were also explored. Through this survey, we found that the carbon containing coated-LiNi0.6Mn0.2Co0.2O2/argyrodite (Li6PS5Cl, synthesized by solid-state method) pair gave the best electrochemical cycling performance with a sustained reversible discharge capacity of 141 mAh.g -1 as well as less than 3% decay upon 20 cycles.

Experimental Section

Solid Electrolyte Synthesis

The solvent-mediated synthesis of β-Li3PS4 solid electrolyte was prepared in tetrahydrofuran (THF).

Stoichiometric amounts of Li2S (Alfa Aesar, purity 99 w/w %) and P2S5 (Acros Organics, purity > 98 w/w %) were dispersed in THF (Sigma Aldrich, 99.9 % v/v). The precursors were stirred continuously for 48 h at room temperature and the powder was later recovered by centrifugation at 600 rpm for 3 min. Next, the obtained powder was placed in a cylindrical Schlenk flask and dried at 100 o C (heating rate 5 o C/min) for 24 h and subsequently at 155 o C for 24 h under dynamic vacuum (P < 0.1mbar).

The solvent-free β-Li3PS4 solid ion conductor was synthesized by mechanochemical milling of a stoichiometric mixture of Li2S and P2S5 using planetary ball-mill followed by annealing step. The mixture of precursors was placed in a ZrO2 jar of 45 mL together with 12 ZrO2 balls (10 mm diameter, 3 g each).

Ball to powder weight ratio was 36 : 1. The mixture was milled at 510 rpm for 20 h applying 15 min milling followed by 15 min of cooling step. The ball-milled powder was later pelletized and annealed at 200 o C (heating rate 5 o C/min) for 1 h in a quartz tube sealed under vacuum. After the annealing process, the sample was cooled naturally.

The solvent-engineered Li6PS5Cl was synthesized in THF/Ethanol solvents. Firstly, a suspension of Li3PS4.xTHF/THF was prepared by stirring of stoichiometric mixture of Li2S and P2S5 in THF at room temperature for 48 h. Next, stoichiometric mixture of Li2S and LiCl (Alfa Aesar, anhydrous, 99 wt. %) was dissolved in ethanol and further stirred at room temperature for 24 h. After the stirring processes, the suspension was transferred to the solution and greenish clear solution was obtained. The final greenish solution was later left stirring continuously overnight and further centrifuged at 600 rpm for 3 min to remove unreacted particles. Finally, the greenish clear solution was dried under dynamic vacuum (P < 1 mbar) to remove the solvent and then, a yellowish powder was obtained. The powder was introduced to a cylindrical Schlenk flask for further drying at 100 o C for 24h and subsequently at 155 o C for 48 h under dynamic vacuum (P < 0.1mbar). Next, the powder was pelletized and annealed at 550 o C (heating rate 5 o C/min) for 6 h in a quartz tube sealed under vacuum. The pelletized sample was cooled naturally after the annealing step.

The solid-state synthesis of Li6PS5Cl was carried out by annealing of stoichiometric mixture of Li2S, P2S5 and LiCl in a Al2O3 crucible. The powder mixture placed in the crucible was sealed under vacuum in a quartz tube and finally, annealed at 550 o C (heating rate 5 o C/min) for 72 h followed by natural cooling.

The water-assisted synthesis of Li3InCl6 was prepared by dissolution of InCl3 (Alfa Aesar, anhydrous, 99,99 w/w %) and LiCl in distilled water. The precursors were left stirring overnight continuously at room temperature and the clear solution was naturally dried at 100 o C. Following the process, white powder was obtained and subsequently it was further dried at 100 o C for 24 h and later at 200 o C for 24 h under dynamic vacuum (P < 1mbar) followed by natural cooling.

The solvent-free Li3InCl6 solid ion conductor was synthesized by mechanochemical milling of a stoichiometric mixture of InCl3 and LiCl using planetary ball-mill followed by annealing step. The ballmill apparatus were used as explained above for solvent-free synthesis of β-Li3PS4 solid ion conductor. The milling process was set at 510 rpm for 24 h applying 15 min milling followed by 15 min of cooling step.

After the ball-milling process, the obtained powder was pelletized and annealed at 260 o C (heating rate 5 o C/min) for 5 h in a quartz tube sealed under vacuum. After the annealing process, the sample was cooled naturally.

Synthesis of core-shell-like cathode composites

β-Li3PS4-coated NMC622 particles were synthesized as follows. A mixture of Li2S (0.115 mg), P2S5 (0.185 mg) and coated-NMC622 (700 mg) was transferred in THF (in a ratio of 35 mg of Li2S and P2S5 solid particles per milliliter of solvent) and stirred for 48 hours at room temperature and subsequently, VGCF particles (50 mg) were added into the suspension by aiming a cathode composite in 70:30:5 wt% (coated-NMC622:β-Li3PS4:VGCF). The solid particles were collected by centrifugation at 600 rpm for 5 min and heated at 100 o C for 24h and then 155 o C for 24 h under vacuum.

Li6PS5Cl-coated NMC622 cathode composite was synthesized in the same line of the solvent-mediated synthesis of Li6PS5Cl solid electrolyte. Firstly, a suspension of Li3PS4.xTHF/THF was obtained at room temperature and next, the obtained suspension was transferred into a solution containing dissolved LiCl and Li2S in ethanol. Next, coated-NMC622 (700 mg) was added into the greenish solution and stirred for 2 h at room temperature. Finally, the solution containing cathode active materials and Li6PS5Cl was dried to remove the solvent at 100 o C for 24h and later at 155 o C for 48 h under vacuum.

Li3InCl6-coated NMC622 particles were obtained using the dissolution of Li3InCl6 particles (300 mg) in THF. After 5 minutes of exposure, coated-NMC622 (700 mg) particles were transferred into the clean solution and stirred for 1 h. Finally, the solution containing cathode active materials and dissolved Li3InCl6 was heated first at 100 o C for 24 h and then at 200 o C for 24 h under vacuum to dry the solvent.

Conductivity measurements

The solid electrolyte samples were pelletized at 4 t/cm 2 during 5 minutes in an 8 mm die set and later transferred to the testing cell for the ionic and electronic conductivity measurements. Ionic conductivity measurements were performed by electrochemical impedance spectroscopy (EIS) at OCV applying excitation amplitude of 50 mV in a frequency range between 7 MHz -1 Hz with 15 points per decade with MTZ impedance analyzer (BioLogic). Electronic conductivity measurements were measured by directcurrent (DC) polarization on the pelletized solid electrolytes with applied voltages of 0.25 V, 0.50 V, 0.75 V and 1 V for 60 min each using a low current channel (BioLogic).

Characterization

Synchrotron X-ray powder diffraction (SXRD) patterns collected at the SOLEIL Synchrotron (St Aubin, France) on the CRISTAL beamline equipped with the Dectris Mythen dectector. All SXRD patterns were collected in transmission mode with λ = 0.58185Å, with the powder sealed in a glass capillary of diameter 0.7 mm. The refinements of the patterns were done using the Rietveld method as implemented in the FullProf program. Laboratory X-ray powder diffraction (XRD) were performed in an airtight cell equipped with a Be window. XRD patterns were recorded in reflection mode in Bragg-Brentano geometry using a Bruker D8 Advance diffractometer equipped with a Cu-Kα X-ray source (λ1 = 1.54056 Å, λ2 = 1.54439 Å) and a LynxEye detector.

Scanning electron microscopy (SEM)

Solid electrolyte and core-shell cathode composite micrographs were obtained on FEI Magellan scanning electron microscope equipped with an Energy dispersive X-ray spectroscopy (EDX) Oxford Instrument detector. EDX of the composite cathode was carried out using an acceleration voltage of 20 kV.

Battery assembling and electrochemical testing

The battery assembly and electrochemical testing were carried out in a cell, which was developed in our laboratory and consisting of a cylindrical polyetherimide (PEI) cell body and two stainless steel pistons of 8 mm diameter.

The two-electrode cell was assembled as follows. 30 mg of solid electrolyte were firstly added into the cell and cold pressed at 1 t/cm 2 for few seconds under a hydraulic press. 16 to 18 mg/cm 2 of cathode composite(coated-NMC622/SE/VGCF, 70:30:5 wt%) was then spread onto one side of the pre-compressed solid electrolyte and next, a mixture of Li0.5In and solid electrolyte (in a 60:40 weight ratio) was added onto other side of the solid electrolyte pellet as counter electrode. Finally, the whole stack was further densified at 4 t/cm 2 for 15 minutes. After the compression, the fully assembled cell was later closed with a torque key applying 2.3 Nm torque in each screw, which yields an internal pressure of 1 t/cm 2 . All assembling procedures were carried out under argon atmosphere in a glove box ([O2] < 1 ppm, [H2O] < 1 ppm).

The three-electrode cell setup was modified from two electrode cell by placing 15 µm Al metal foil as RE current collector between two PEI body parts and the cell assembly was realized as explained in our previous work. [START_REF] Dugas | Engineered Three-Electrode Cells for Improving Solid State Batteries[END_REF] Galvanostatic cycling studies were carried out in the voltage range of 2.1 -3.6 V vs. LiIn/In at room temperature at C/20 (C correspond to 1 mole of Li per mole of active material in 1 h). Three-electrode fullcell impedance spectra were measured upon galvanostatic cycling at 3.2 V, 3.4 V, 3.6 V, 3.4 V, 3.2 V and 2.1 V after 4 h rest at each point using a 10 mV sinusoidal perturbation. 10 points per decade were taken for frequencies between 1 MHz -10 mHz.

All electrochemical measurements were conducted with a VMP3 potentiostat/galvanostat (BioLogic) controlled with EC-Lab software.

Results and Discussion

Pure solid electrolytes β-Li3PS4, Li6PS5Cl and Li3InCl6 powders were synthesized via solvent-assisted and solvent-free methods according to experimental protocols derived from the literature (see experimental section) and will be herein denoted with the prefixes W-and C-, respectively (Figure 1). The solid electrolytes were first characterized for their structure, morphology and electronic-ionic conduction by XRD, SEM and EIS spectroscopy, respectively. The collected XRD patterns (Figure 2a, b andc) match those already reported and indicate that nearly all the samples are pure phases. As expected, a greater broadening of the peaks was observed for samples made by mechanical ball milling given that this route promotes microstructural defects and particle nano-sizing. It is to note that Li6PS5Cl can be prepared as single phase by solid-state reaction, as opposed to the solution route, which also yields several crystalline impurities (see symbols in figure 2b) identified as Li3PO4, Li2S and LiCl (Figure S1). Such a difficulty was also mentioned in the literature and associated to the interaction between PS4 3-unit and ethanol. [START_REF] Zhou | Solvent-Engineered Design of Argyrodite Li 6 PS 5 X (X = Cl, Br, I) Solid Electrolytes with High Ionic Conductivity[END_REF][START_REF] Ruhl | Impact of Solvent Treatment of the Superionic Argyrodite Li 6 PS 5 Cl on Solid-State Battery Performance[END_REF] Figure 1: Synthesis routes and abbreviations used for studied solid electrolytes.

Figure 2d-i shows the SEM micrograph of the synthesized powders whose particles adopt various shapes depending upon the electrolyte chemistry. Particle morphology is nearly independent of the synthesis strategy for the halides in contrast to the noticeable differences for β-Li3PS4 and Li6PS5Cl when prepared by solid-state and wet route. The relative density of pelletized solid electrolytes (SEs) was also calculated, revealing that solvent-free route seems to facilitate the material densification enabling a porosity decrease down to ~5% for Li3InCl6 (Figure S2). Room temperature ionic and electronic conductivities were determined for all the SEs by EIS and DC polarization (Figure S3), respectively, and are shown in Figure 3a and Figure 3b together with the corresponding EIS spectra in Figure 3c, d ande. A slightly better ionic conductivity was obtained for the solvent-free electrolytes compared to their counterpart prepared by wet route with the only exception of W-Li3InCl6. The obtained ionic conductivities rank as follows: Li6PS5Cl (4x10 -3 S.cm -1 ), β-Li3PS4 (0.66x10 -3 S.cm -1 ) and Li3InCl6, (0.64x10 -3 S.cm -1 ) and are in line with previous reports. Turning to the DC electronic conductivity, it remains nearly constant and equal to 10 -9 -10 -10 S.cm -1 for most of SEs except for the W-Li6PS5Cl whose electronic conductivity is three orders of magnitude higher (1.1x10 -6 S.cm -1 ) and exceed the values reported in the literature. [START_REF] Zhou | Solvent-Engineered Design of Argyrodite Li 6 PS 5 X (X = Cl, Br, I) Solid Electrolytes with High Ionic Conductivity[END_REF] Interestingly, the DC electronic conductivity determined for the same W-Li6PS5Cl prior to the firing at 550°C was 1.8x10 -9 S.cm -1 (Figure S3g). Thus, we believe that the enhanced electronic conductivity of W-Li6PS5Cl after firing could be rooted in the pyrolysis of residual solvent at high temperature (550 o C) yielding carbonaceous species. Overall, the highest ionic and lowest electronic conductivities were achieved by C-Li6PS5Cl and W-Li3InCl6 electrolytes.

Cathode composites and battery testing

Following the characterization of the pure solid electrolytes presented in the previous section, we proceeded to pair them with a high voltage cathode material and test the performance of the resulting composite in full ASSBs. To this effect, ASSBs containing the solid electrolytes under study were assembled and tested as described in the experimental section. A single positive electrode material coated-LiNi0.6Co0.2Mn0.2O2

(referred to as NMC622) was used throughout this work. In the case of β-Li3PS4 and Li6PS5Cl, the following general configuration was used: coated-LiNi0.6Co0.2Mn0.2O2/SE/VGCF | SE | Li0.5In/SE, where SE refers to the corresponding electrolyte. However, such configuration was not suitable to test Li3InCl6 due to its known instability at the anode interface leading to the reduction of In 3+ . [START_REF] Riegger | Lithium-Metal Anode Instability of the Superionic Halide Solid Electrolytes and the Implications for Solid-State Batteries[END_REF] Thus, a hetero-structure cell design was employed. It consists in the use of Li3InCl6 only in the cathode composite, C-Li6PS5Cl as separator and a composite made of C-Li6PS5Cl and Li0.5In as the anode. At least three cells were cycled for each cathode composite in order to assess reproducibility of the results.

Pairing coated-NMC622 with electrolytes obtained from solvent-free route (ceramic)

The voltage profiles of the first galvanostatic charge/discharge cycle as well as the evolution upon cycling of the discharge capacity are comparatively shown, respectively, in Figures 4a,4b for the full-batteries containing electrolytes prepared by solvent-free route. Additionally, the cell polarization (ΔV, difference between average charge and discharge voltages as shown in Figure S4a) is depicted in Figure 4c.

The cell using coated-NMC622/C-Li6PS5Cl/VGCF composite positive electrode showed initial charge and discharge capacities of 173 and 141 mAh.g -1 , respectively (Figure 4a) with a 81.5% of Coulombic efficiency and a remarkable high capacity retention of 97% up to 20 cycles. (Figure 4b). With such figures of merits, C-Li6PS5Cl-based cells outpaces the ones containing C-Li3InCl6 and C-β-Li3PS4. In contrast, the analogous cell with a carbon-free cathode composite exhibited a limited initial capacity of 73 mAh.g -1 , which gradually increased up to 101 mAh.g -1 after 20 cycles (Figure S5a and5b), most likely due to progressive activation of the CAM particles upon cycling. Additionally, such carbon-free batteries showed a higher cell resistance compared to their carbon-containing counterparts (Figure S6), in line with its poorer cycling performance. Overall, this indicates that the addition of carbon to the composite containing C-Li6PS5Cl enhanced the reversible capacity of the battery and its overall cycling performance without any major degradation (Figures 4c).

Turning to the cell based on carbon-containing coated-NMC622/C-β-Li3PS4 composite, it showed an initial charge and discharge capacities of 156 and 109 mAh.g -1 , respectively (Figure 4a) with a 10% capacity decay after 20 cycles (Figure 4b) as opposed to the carbon-free composite, which exhibited an almost constant discharge capacity of 110 mAh.g -1 (Figure S5b). Moreover, the carbon-containing composite showed an increase in polarization ΔV upon cycling (Figures 4c). In summary, the addition of carbon to β-Li3PS4-based composite electrodes resulted in a capacity boost but at expenses of a poor Coulombic efficiency and further capacity fading. Such a negative effect of carbon has been ascribed to its electron mediator role to trigger electrochemical-driven degradation. [START_REF] Li | Electrochemically Primed Functional Redox Mediator Generator from the Decomposition of Solid State Electrolyte[END_REF] The greater deterioration of β-Li3PS4 -carbon composites rather than Li6PS5Cl -carbon ones is rooted in the lower stability against deterioration of the former at oxidation potentials greater than 3.2 V as deduced independently by EIS experiments.

Lastly, carbon-containing composite electrodes with C-Li3InCl6 (coated-NMC622/C-Li3InCl6) was tested using C-Li6PS5Cl as separator. High charge and discharge capacities were delivered in the first cycle (166 and 143 mAh.g -1 , respectively, Figure 4a), comparable to our best performing material (C-Li6PS5Cl-based composite). However, its capacity fading upon cycling resulted the worst of the three investigated systems, reaching 50% after 20 cycles (Figure 4b). Such a rapid capacity decay is mirrored by a fourfold increase of ΔV upon cycling (Figure 4c). Moreover, when replacing C-Li6PS5Cl testing separator by W-β-Li3PS4 separator, we observed an even greater capacity fade that reached 95% after cycling (Figure S7). In order to interrogate the origin of such a drastic capacity fading, we investigated the chemical compatibility between W-Li3InCl6 and the aforementioned electrolytes also used in the cell architecture. To this effect, C-Li3InCl6/C-Li6PS5Cl and C-Li3InCl6/W-β-Li3PS4 mixtures of the pristine powders (in 1:1 wt%) were prepared and subjected to a thermal treatment at 150°C during 72 hours in order to promote chemical reactivity. Thereafter, synchrotron X-ray powder diffraction of the annealed powders revealed the presence of LiCl as the decomposition product, thus confirming thermally promoted chemical reactions between the two solid electrolytes (Figure S8a andS8b). This degradation was certainly more pronounced in presence of b-Li3PS4, in consistence with the rapid performance degradation of halide-based cells using b-Li3PS4 as the separator. Bearing in mind that our heating treatment was below the synthesis temperature of the mother electrolytes (W-β-Li3PS4: 155 °C, C-Li6PS5Cl: 550 °C and C-Li3InCl6: 200 o C) and therefore far from their unstable thermal stability range, we can relate the newly found chemical instability of Li3InCl6 towards β-Li3PS4 and Li6PS5Cl to the cycling performance degradation of the hetero-structural battery. Such chemical incompatibility leads to the formation of a passivation interlayer between the positive electrode and the separator that is prejudicial to battery performances. Inspired by earlier literature reports, we have replaced our hand grinding-based procedure for cathode composite preparation by a process consisting in seeding our reacting electrolyte solution in presence of cathode active material powders (coated-NMC622). The idea behind such new strategy is to produce cathode particles coated with SE in a core-shell like structure and thus to favor the intimate contact between these two components. However, such an implementation was not free of complications. Synthesis limitations became quickly evident for the argyrodite-based and halide-based composites, namely during the subsequent temperature treatment of the solution. For instance, we experienced a structural degradation of W-Li6PS5Cl-coated NMC622 powders when firing at 550 °C as shown by XRD (Figure S15). We thus limited the annealing temperature to 155 °C thus sacrifying the SE ionic conductivity which reaches its maximum (3.5x10 -3 S.cm -1 ) for W-Li6PS5Cl powders fired at 550 °C. Moving to halides did not simplify the task since re-crystallized W-Li3InCl6 from THF solution at 200 °C in either absence or presence of coated-NMC622 particles revealed a multi-phased product as evidenced by XRD (Figure S16a) that has an order of magnitude lower ionic conductivity (Figure S16b). In contrast, W-β-Li3PS4-coated NMC622 based composites could be successfully prepared at 155 °C via our solution process without any major roadblock.

Figures 6a-c show SEM micrographs and EDX analysis for the prepared composites based on W-β-Li3PS4, W-Li6PS5Cl and W-Li3InCl6. In all cases, the rounded rather than salient particle edges pertaining to coated-NMC622 suggest the presence of deposits. However, they differ in their EDX mapping indicating a uniform repartition of Ni and P for the W-β-Li3PS4-coated NMC622 particles as opposed to a non-uniform repartition for the two other composites (see Cl and In mapping on Figure 6b and6c, respectively), hence confirming the difficulties in achieving efficient core-shell coated particles. Despite these differences, the cycling performance of the composites were tested in full ASSBs assembled within the same electrode-electrolyte configuration and results are solely reported after duplicated measurements. The results are recapped in Figures 6d-f, respectively. The composite based on W-Li3InCl6 showed the worst performance, and this does not come as a surprise in light of phase purity and non-uniform deposit issues deduced from combined XRD and conductivity results. Similar but less pronounced problems rendered the performance of W-Li6PS5Cl-coated composites, which resulted worse capacity than those obtained previously on hand-ground Li6PS5Cl-based cathode composites. Turning to the performance of W-β-Li3PS4-coated NMC622 cathode composite, it nearly compares to that of their hand-ground counterpart with in one way a lower reversible capacity (118 vs. 129 mAh.g -1 ) but a slightly higher capacity retention after 20 cycles (78% vs. 75.8%). It is worth mentioning that adding carbon additives to this composite electrode resulted in a deterioration of the capacity retention, in line of what observed before (Figure S17). Altogether, these results provide no advantage of the in-situ solution coating of coated-NMC622 particles by SEs. Other approaches towards the synthesis of tailor made core-shell particles remains to be developed. Li6PS5Cl or β-Li3PS4 that translates into the formation of a passivation interlayer in the hetero-structure cell design, which leads to degraded electrochemical performance up to 50% upon cycling. Obviously, the presence of a triple phase NMC-Li6PS5Cl-Li3InCl6 contact point within a composite is detrimental. Thus, designing hetero-structured separators together with anodic protective approaches stands as common strategies to overcome these interface issues. They are being explored in our group and preliminary data indicate promising results, hence enabling to envision the future use of halides in solid-state batteries.

Overall, we hope this work could provide guidance in properly selecting optimized electrolyte-electrode pairs for developing practical all-solid-state batteries.

Associated Content

Supporting Information

Figure 2 :

 2 Figure 2: Powder X-ray diffraction patterns of a-c) W-β-Li3PS4 (dark green), C-β-Li3PS4 (green), W-Li6PS5Cl (wine), C-Li6PS5Cl (red), W-Li3InCl6 (navy) and C-Li3InCl6 (blue). Asterisks indicate Be reflections from the sample holder. SEM images of d) W-β-Li3PS4, e) C-β-Li3PS4, f) W-Li6PS5Cl, g) C-Li6PS5Cl, h) W-Li3InCl6 and i) C-Li3InCl6 as synthesized.

Figure 3 :

 3 Figure 3: Comparison of a) ionic and b) electronic conductivities of studied solid electrolytes and Nyquist plots of c) β-Li3PS4, d) Li6PS5Cl and e) Li3InCl6 prepared through solvent-engineered and solvent-free methods.

Figure 4 :

 4 Figure 4: Electrochemical performance results of hand-ground carbon-containing coated-NMC622 and β-Li3PS4 / Li6PS5Cl / Li3InCl6 cathode composites in ASSBs at room temperature. Voltage profile of the first galvanostatic charge-discharge cycle, discharge capacity as a function of cycle number and cell polarization over cycling depicted, respectively, for the cells containing electrolytes prepared by solvent-free (a, b, c) and wet-based (d, e, f) synthesis routes.

Figure 5 :

 5 Figure 5: Galvanostatic voltage profiles of W-Li3InCl6/coated-NMC622/VGCF positive electrode, C-Li6PS5Cl as separator and Li0.5In/C-Li6PS5Cl negative electrode in three-electrode cell configuration a) 1 st and 20 th cycle curves. The impedance spectrums of b) the W-Li3InCl6/coated-NMC622/VGCF positive electrode, where the symbols show first and last impedance sequence over cycle and the black arrows indicate the evolution of the impedance at each cycle (offset was added between each impedance spectrums) and c) the Li0.5In/C-Li6PS5Cl negative electrode for each charge and discharge sequences for 1 st and 20 th cycles.

  with solvent-based electrolytes in a core-shell architecture

Figure 6 :

 6 Figure 6: SEM images and the corresponding EDX elemental maps (representative elements nickel, phosphorous, chlorine and indium are shown in yellow, turquoise, blue and red for coated-NMC622, W-β-Li3PS4, W-Li6PS5Cl and W-Li3InCl6, respectively) of a) W-β-Li3PS4-coated NMC622, b) W-Li6PS5Cl-

Figure 7 :

 7 Figure 7: Comparison of the charge (dark colored bars) -discharge (light colored bars) capacities and the cell polarization of the all studied carbon-based cathode composites between 1 st and 20 th cycles.

Acknowledgements

The authors thank Bernhard Leube for fruitful discussions on argyrodite synthesis, Marie-Francine Lagadec for EDX imaging, Biao Li and Parth Desai for helpful discussions. We acknowledge the precious help of Benoît Baptiste (IMPMC, Sorbonne Université) who collected the SXRD patterns at the CRISTAL beamline of Synchrotron Soleil (BAG proposal number 20201440); staff from SOLEIL is also acknowledged.

Pairing coated-NMC622 with electrolytes obtained from wet route

The initial charge/discharge performances and the discharge capacity upon cycling are shown in Figures 4d and4e respectively, for the full-batteries containing solid electrolytes synthesized by wet route.

Additionally, the ΔV (the difference between average charge and discharge voltages as shown in Figure S4b) as a function of the cycle number is depicted in Figure 4f.

The cathode composite of the W-Li6PS5Cl (coated-NMC622/W-Li6PS5Cl/VGCF) exhibited a remarkably high first charge capacity of 246 mAh.g -1 accompanied by a pronounced oxidation slope starting at ~ 2.3 V, likely ascribed to electrolyte decomposition. In line with such hypothesis, the Coulombic efficiency of the first cycle was very poor (54.3%) (Figure 4d). This suggests a magnified electrochemical decomposition for the W-Li6PS5Cl, which may be possibly related to faster decomposition reaction [START_REF] Nakamura | Guidelines for All-Solid-State Battery Design and Electrode Buffer Layers Based on Chemical Potential Profile Calculation[END_REF] triggered by its unexpected very high electronic conductivity shown in the previous section together. Nevertheless, after completing the first cycle, the system showed a striking 96.8% of capacity retention and an almost constant Coulombic efficiency of 86.5% after 20 cycles (Figures 4e andS9, respectively). Note that, variation on fabrication conditions of W-Li6PS5Cl did not deliver better cycling performance (see Figure S10). Equally, the system showed a nearly stable ΔV and apparent resistance upon cycling represented in Figures 4f andS11, respectively. Interestingly, the carbon-free positive electrode displayed an increasing discharge capacity upon battery cycling starting from 67 mAh.g -1 up to 100 mAh.g -1 (Figure S5e), possibly related to a gradual activation of the CAM given along cycling, in the line with C-Li6PS5Cl reported above. Such an activation period upon cycling is nearly common to all the carbon-free positive electrodes. A possible reason could lie the repetitive contraction-expansion of the active material that enables through cycling a better contact between the active material and solid electrolyte, hence enabling a greater participation of the NMC particles and by the same token an increase in capacity.

When moving to the W-β-Li3PS4, carbon-containing composite (coated-NMC622/W-β-Li3PS4) exhibited a relatively high charge and discharge capacities upon the first cycle (171 and 129 mAh.g -1 , respectively) that is accompanied by a rapid capacity decay of 71.5% after 20 cycles. This differs from the analogous carbonfree composite, which delivered a stable discharge capacity of 58 mAh.g -1 upon cycling (Figures S5d andS5e). The differences between the carbon-containing and carbon-free composites are also reflected in the evolution of the cell polarization (DV) which either increases (Figure 4f) or remains constant (Figure S5f) upon cycling, respectively. We believe that such a carbon-driven decrease in performance is nested in the electrochemical degradation of the solid electrolyte that is evidenced as a sloping region at the early stage of the first charge (~2.3V), in line with the reported study in literature. [START_REF] Tan | From Nanoscale Interface Characterization to Sustainable Energy Storage Using All-Solid-State Batteries[END_REF] Aside from such detrimental effects, a beneficial effect of carbon is that it enables to tap more capacity from the positive composite electrode via improved electronic conductivity as already well-described. [START_REF] Tan | From Nanoscale Interface Characterization to Sustainable Energy Storage Using All-Solid-State Batteries[END_REF][START_REF] Walther | Influence of Carbon Additives on the Decomposition Pathways in Cathodes of Lithium Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Zhang | The Detrimental Effects of Carbon Additives in Li 10 GeP 2 S 12 -Based Solid-State Batteries[END_REF] Lastly, we tested W-Li3InCl6 in a dual solid electrolyte cell architecture, as described above. A first charge and discharge capacities of 165 and 142 mAh.g -1 respectively (Figure 4d) was obtained together with the lowest irreversible capacity through the first cycle (13.6%) obtained among all the cathode-solid ionic conductor pairs tested in this work. However, a progressive capacity degradation was observed, reaching a 50% of decay after 20 cycles (Figure 4e) as witnessed by an increase in ΔV (Figures 4f), analogous to the case of C-Li3InCl6 reported above. As it was previously shown, the scenario upon cycling became worse when switching to W-β-Li3PS4 as separator, with a rapid capacity decay (11.7% after 20 cycles) (Figure S12). To further evaluate the origin of the electrochemical degradation in the dual solid electrolyte cell structure, W-Li3InCl6-based cathode composite was tested in a three-electrode cell configuration using C-Li6PS5Cl as separator and Li0.5In/C-Li6PS5Cl as reference and negative electrode. Galvanostatic cycling together with EIS measurements were performed in order to monitor the impedance evolution upon cycling of the positive and negative electrode separately (see experimental details). Figure 5a displayed the comparison of galvanostatic cycling profiles of positive and negative electrodes between first and 20 th cycles, which exhibit a pronounced capacity degradation, consistent with the two-electrode cell setup, as well as higher polarization in positive electrode rather than negative electrode after 20 cycles. On the other hand, the impedance response revealed that the most significant change came from the positive electrode by delivering up to a six-fold increase of the resistance at f= 1Hz after 20 cycles (Figure 5b). Such impedance response was mainly originated in the mid frequency region as opposed to the high frequency region. In contrast to the remarkable impedance growth at the positive electrode interface, no significant impedance enhancement was observed for the negative electrode over cycling (Figure 5c). When reconstructing the hetero-structure with W-β-Li3PS4 as separator, the impedance increase in the positive electrode became even higher over cycling, aligned with its rapid electrochemical cell fading (Figure S13).

Overall, the electrochemical performance degradation in the hetero-structure cell of W-Li3InCl6-based cathode composite upon cycling was found via three-electrode cell-setup to originate from highly resistive positive electrode interface. This interfacial resistance as well as the poor performances of halide-based cells are rooted in the chemical incompatibility between Li3InCl6 towards the two sulfide-based electrolytes used in the heterogeneous cell architecture, evidenced in their thermally driven reactivity at low temperature as shown in Figure S14a andb.

Conclusion

We have herein reported systematic crossed-comparisons of the chemical/electrochemical compatibility of coated-LiNi0.6Mn0.2Co0.2O2 cathode material against the today's three most studied solid ionic conductors prepared via various synthesis approaches. From this survey, which has not taken into accounts effects of particle size and morphology, we found that solvent-free synthesis leads to high ionic and low electronic conduction, impurity-free structure and high densification for solid ionic conductors (except for Li3InCl6).

Furthermore, amongst all studied cathode composites whose electrochemical performance results are recapped in Figure 7, C-Li6PS5Cl and coated-LiNi0.6Mn0.2Co0.2O2 composite positive electrode (handground) was found to enable the best electrochemical cycling performance with a sustained reversible capacity of 141 mAh.g -1 for at least 20 cycles in all-solid-state batteries and preserved the best performance upon 50 cycles as shown in Figure S18. This remarkable performance was linked to the presence of carbon additive to boost the electronic conductivity of cathode composite at expenses of its detrimental effect on sulfide-based solid ionic conductors, as we obtained in β-Li3PS4-based cathode composites with capacity degradation up to 30%. This illustrates two sides of the same coin when considering how adding carbon in the cathode composite will affect the battery performance. Whereas a boost in the electronic conduction of the cathode may be beneficial capacity-wise, it can also promote electrolyte decomposition as deduced from XRD measurements, which will translate into poor Coulombic efficiency and capacity fading. The extent of this decomposition depends upon the thermodynamic stability of the ionic conductor with respect to the active material for which DFT can inform us. A way to alleviate such deterioration of the SE is to use sulfide-based composite electrodes such as Li1.13Ti0.57Fe0.33S2 rather than oxide ones (NMC) which operating voltage falls within the range of the electrolyte stability so that coatings are not needed. Sulfides offers better ionic and electronic conductivities, both leading to a smaller polarization with however a penalty in terms of energy density owing to their lower redox voltage as compared to oxides.

Through this study, we have revealed the limitations of solvent-assisted synthesis approach of solid electrolytes, namely residual solvent and phase purity issues, which perturb both chemical and electrochemical properties. Although conceptually elegant, the in-situ design of core-shell cathode active material-solid electrolyte structures in solution was shown to suffer from detrimental effect of solvents on solid electrolytes or chemical instability between layered oxide and sulfide-based solid electrolyte at high temperature, hence they observed limited performances compared to hand-ground cathode composites.

Halide-based solid electrolytes are particularly attractive for their high voltage stability and therefore for its electrochemical compatibility with oxide-based cathode materials. However, their instability towards Li-based anodes rules out the possibility of having full halide-based cell architectures. To contour this galvanostatic cycling result of W-β-Li3PS4-coated NMC622 (coated) cathode composite (core-shell process) in carbon-free system, electrochemical performance results of carbon-containing cathode composites over 50 cycles.

Notes

The authors declare no competing financial interest.

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