domain: spi chim phys

text:
Wen Lu

Jie Xu

Qingsong Song

Kaiting Bian

Jun Guo

Jian Liu

Dongzhen Li

Peng Liu

Chaojin Zhang

Xiaodong Xu

Jie Xu

Jun Xu
email: xujun@mail.shcnc.ac.cn

Kheirreddine Lebbou
email: kheirreddine.lebbou@univ-lyon1.fr

Pr 3+ deactivation effect to Tm 3+ at ~ 1.45 µm emission in Bi 4 Ge 3 O 12 crystals grown by the micro-pulling-down method

Keywords: Tm, Pr:Bi4Ge3O12, Micro-pulling-down method, Spectroscopic property

Pr3+ deactivation effect to Tm3+ at ∼1.5 µm emission in Bi4Ge3O12 crystals grown by the micro-pulling-down method

Introduction

The near-infrared (NIR) emissions around 1.5 μm wavelength have attracted much attention in the past decades for its numerous applications in telecommunication, biomedical systems and remote sensing [START_REF] Loiko | Emission properties of Tm 3+ -doped CaF2, KY3F10, LiYF4, LiLuF4 and BaY2F8 crystals at 1.5 µm and 2.3 µm[END_REF][START_REF] Diening | Diode-pumped continuous-wave, Quasi-continuous-wave, and Q-switched laser operation of Yb 3+ ,Tm 3+ : YLiF4 at 1[END_REF][START_REF] Allen | An Efficient 1.46µm Thulium Fiber Laser via a Cascade Process[END_REF]. Tm 3+ ion is a natural candidate for 1.5 μm laser owing to the 3 H4 → 3 F4 transition. Unfortunately, the self-terminating transition ( 3 H4 → 3 F4) of Tm 3+ ions, namely, the upper level ( 3 H4), has a shorter lifetime than the lower one ( 3 F4), almost extinguishing the probability of population inversion for lasing at around the 1.5 μm wavelength. In order to achieve intense Tm 3+ : 3 H4 → 3 F4 NIR emissions for practical laser operation, co-doping deactivating ions, such as Tb 3+ [START_REF] Ermeneux | Growth and fluorescence properties of Tm 3+ doped YVO4 and Y2O3 single crystals[END_REF][START_REF] Ryba-Romanowski | Relaxation dynamics of excited states of Tm 3+ in SrGdGa3O7 crystals activated with Tm 3+ and Tm 3+ +Tb 3+[END_REF][START_REF] Zhang | Enhanced 1.4 µm emissions of Tm 3+ via Tb 3+ deactivation in (Gd0.5Lu0.5)2SiO5 crystal[END_REF][START_REF] Li | Growth and spectroscopic properties of Tm 3+ and Tb 3+ co-doped GdScO3 crystal[END_REF][START_REF] Bettinelli | Fluorescence dynamics of YVO4:Tm 3+ , YVO4:Tm 3+ ,Tb 3+ and YVO4:Tm 3+ ,Ho 3+ crystals[END_REF][START_REF] Braud | Spectroscopy and Fluorescence Dynamics of (Tm 3+ , Tb 3+ ) and (Tm 3+ , Eu 3+ ) Doped LiYF4 Single Crystals for 1.5 µm Laser Operation[END_REF], Ho 3+ [START_REF] Bettinelli | Fluorescence dynamics of YVO4:Tm 3+ , YVO4:Tm 3+ ,Tb 3+ and YVO4:Tm 3+ ,Ho 3+ crystals[END_REF]and Eu 3+ [START_REF] Braud | Spectroscopy and Fluorescence Dynamics of (Tm 3+ , Tb 3+ ) and (Tm 3+ , Eu 3+ ) Doped LiYF4 Single Crystals for 1.5 µm Laser Operation[END_REF][START_REF] Lisiecki | Relaxation of excited states of Tm 3+ and Tm 3+_ Eu 3+ energy transfer in YVO4 crystal[END_REF], is an effective way to depopulate the Tm 3+ : 3 F4 level for population inversion and create a desirable fluorescence lifetime ratio. Moreover, Tm 3+ lasers emitting at around 1.5 μm have been demonstrated in Tm 3+ /Ho 3+ co-doped YLiF4 crystal [START_REF] Stoneman | Continuous-wave 1.50-µm thulium cascade laser[END_REF], Tm 3+ /Tb 3+ co-doped YLiF4 crystal [START_REF] Rosenblatt | Diode-Pumped Room-Temperature cw 1.45-µm Tm,Tb:YLF laser, Advanced Solid State Lasers[END_REF] and Tm 3+ /Ho 3+ co-doped tellurite glass microsphere [START_REF] Li | Tm 3+ -Ho 3+ codoped tellurite glass microsphere laser in the 1.47 µm wavelength region[END_REF]. To the best of our knowledge, there is no report about the use of Pr 3+ as a deactivated ion of Tm 3+ for achieving an efficient 1.5 μm fluorescence emission. Bi4Ge3O12 (BGO) crystal has some excellent characters as a host of laser gain medium, such as excellent physical and chemical properties [START_REF] Williams | Optical, thermo-optic, electro-optic, and photoelastic properties of bismuth germanate (Bi4Ge3O12)[END_REF]. Undoped BGO has broad emission bands which can be converted to narrow bands emission lines when it is doped with rare earth. The crystal belongs to cubic system with the space group of I43d [START_REF] Farias | The local structure around the Nd impurity incorporated into the Bi4Ge3O12 crystal matrix: An ab initio study[END_REF]. As it melts congruently, large size single crystals can be obtained by the Bridgman and Czochralski methods. The micro-pulling-down (µ-PD) method can also be used to grow fiber shaped BGO crystals with high quality and good transparency, where the crystals have low thermal strain compared to others growth methods [START_REF] Shim | Growth of Bi4Ge3O12 single crystal by the micro-pulling-down method from bismuth rich composition[END_REF][START_REF] Chani | Evaporation induced diameter control in fiber crystal growth by micro-pulling-down technique: Bi4Ge3O12[END_REF][START_REF] Farhi | BGO fibers growth by μ-pulling down technique and study of light propagation[END_REF]. Much research work has been devoted to Nd 3+ [START_REF] Li | Spectroscopic properties of Eu:Bi4Ge3O12 single crystal grown by the micro-pulling down method[END_REF][START_REF] Lin | Growth and laser properties of Nd 3+ -doped Bi4Ge3O12 single-crystal fiber[END_REF], Eu 3+ [START_REF] Shim | Eu 3+ doped Bi4Ge3O12 fiber crystals grown by the by the micro-pulling down method[END_REF][START_REF] Li | Optical properties of Nd:Bi4Ge3O12 crystals grown by the micro-pulling down method[END_REF] doped BGO crystals grown by the µ-PD method as laser gain media. However, Tm 3+ doped BGO crystal for NIR lasing emission has not been reported yet.

In this work, Tm 3+ /Pr 3+ co-doped BGO shaped crystals were successfully grown by the µ-PD method. Pr 3+ was demonstrated to be an efficient deactivated ion for Tm 3+ ion to facilitate the Tm 3+ : 3 H4 → 3 F4 by effective energy transfer from Tm 3+ : 3 F4 level to Pr 3+ : 3 F2 level. Pr 3+ and Tm 3+ are attractive dopants because they have many emission lines due to 4f-4f transitions from visible to near infrared wavelengths. The spectroscopic properties of Tm 3+ /Pr 3+ co-doped BGO crystals at the 1.5 μm wavelength region for potential NIR laser application was also evaluated. Following our knowledge, not any work with Pr, Tm codoping BGO crystals study luminescence spectra in the NIR region has been investigated.

Experiments

The Tm,Pr:BGO crystals were grown using the µ-PD method [START_REF] Song | Structure and white LED properties of Ce-doped YAG-Al2O3 eutectics grown by the micro-pulling-down method[END_REF] with the following composition: Tm0.08Pr4xBi3.92-4xGe4O12 (x=0, 0.001, 0.002, 0.003, 0.005). Tm2O3 (99.999%), Pr6O11 (99.999%), Bi2O3 (99.999%) and GeO2 (99.999%) powders were used as raw materials for crystal growth. The mixture was press into bulks and heated in the air at 750℃ for 24h. The polycrystalline powders was melted in a platinum crucible, and then the melt was pulled down along the <210> direction through a capillary channel at the bottom of the crucible. The crystals were grown with a pulling rate of 0.3 mm/min in the air atmosphere. As shown in Fig. 1, Tm,Pr:BGO cubic structure crystals were obtained with a rough appearance on the outside and present a light yellow color typical to BGO crystal growth from the melt in air atmosphere using Pt crucible. The crystals crystallized from the capillary die of the crucibles were very uniform in diameter according to the measurements performed directly from the fibers with accuracy of about ± 20 μm. The absorption spectra of Tm,Pr:BGO crystals in the range of 400-2050 nm were measured using a UV-VIS-NIR spectrophotometer (Lambda 1050, Perkins Elmer, U. S. A.). The fluorescence spectra and fluorescence decay curves were recorded by a fluorescence spectrometer (Edinburg Instrument, FLS1000, U. K.) under the excitation of 793 nm. All the measurements were performed at room temperature.

Results and discussion

The typical absorption spectra of Tm:BGO and Tm,Pr:BGO crystals are presented in Fig. 2. The observed absorption bands are assigned to the transitions from the ground states of Tm 3+ : 3 H6 or Pr 3+ : 3 H4 to the indexed excited ones. Whatever the dopant concentration, no spectral shifts of the absorption spectra were detected. Intensities of the main absorption are mostly similar at room temperature (RT), but relative changes of line intensities are observed. In the range of 760-820 nm, there is one strong absorption band near 792 nm with a full width at half-maximum (FWHM) about 3 nm due to the Tm 3+ : 3 H6 → 3 H4 transition, which coincides well with the emitting wavelength of high-power AlGaAs laser diodes. Since Pr 3+ ions do not absorb around 793 nm, they will not reduce the pump efficiency under laser diode excitation. 3(a) and Fig. 3(b). The NIR emission bands correspond to the Tm 3+ : 3 H4 → 3 F4 transition and The MIR emission bands are assigned to the Tm 3+ : 3 F4 → 3 H6 transition. It is clear to see that almost no distinct change in the shape and position of the NIR emission bands were observed at praseodymium concentrations from 0 to 0.5 at.%. The presence of the narrow emissions lines resulting from dopant transition offers a probe of the geometry and crystal field. Moreover, the intensity of the emission bands changed a little. On the contrary, a significant reduction in the MIR emission intensity of the 3 F4 level between Tm 3+ doped and Tm 3+ /Pr 3+ co-doped samples, which justifies that Pr 3+ ions can be used effectively to depopulate the Tm 3 +: 3 F4 level. The intensity of MIR emission decreases with the increase of praseodymium concentration, and the intensity of MIR emission is almost absent when the Pr 3+ concentration is over 0.3 at.%. In order to understand the differences of fluorescence spectra caused by co-doping Pr 3+ , the energy level scheme of Tm 3+ and Pr 3+ is illustrated in Fig. 4. The Tm 3+ ions were excited from ground state 3 H6 to excited state 3 H4 through the absorption of 793 nm pumping, and then the ions in the 3 H4 level decay radiatively to the 3 F4 level with 1.5 µm emission. After that the Tm 3+ ions in the 3 F4 level will decay radiatively to the 3 H6 level with emitting 2 µm emission. On the other hand, energy transfer process from 3 F4 level of Tm 3+ to Pr 3+ : 3 F4 (Tm 3+ ) + 3 H4 (Pr 3+ ) → 3 H6 (Tm 3+ ) + 3 F2 (Pr 3+ ) will also happen because of the small energy mismatch between the 3 F2 level of Pr 3+ and the 3 F4 level of Tm 3+ , then the energy will decay very rapidly down to the ground state via multiphonon relaxation. This process would depopulate the 3 F4 level of Tm 3+ and quenching the MIR emission of Tm 3+ . ET: energy transfer from Tm 3+ : 3 F4 level to Pr 3+ : 3 F2 level.

In order to further confirm the energy interaction mechanism, the fluorescence decay curves of Tm 3+ : 3 H4 and 3 F4 multiplets of Tm:BGO and Tm,Pr:BGO crystals were measured at 1452 and 1879 nm, as shown in Fig. 5. and Fig. 6. The measured decay curves were fitted with a single exponential function and the fluorescence lifetimes τ are listed in Table 1. It is noticed that the fluorescence lifetime of the 3 H4 state (218 µs) is much shorter than that of 3 F4 state (5.64 ms) in Tm:BGO crystal due to the cross-relaxation process: 3 H4 (Tm 3+ ) + 3 H6 (Tm 3+ ) → 3 F4 (Tm 3+ ) + 3 F4 (Tm 3+ ) [START_REF] Braud | Spectroscopy and Fluorescence Dynamics of (Tm 3+ , Tb 3+ ) and (Tm 3+ , Eu 3+ ) Doped LiYF4 Single Crystals for 1.5 µm Laser Operation[END_REF]. As we know, the lifetime of terminal level 3 F4 is too long compared with that of the emitting level 3 H4, which can cause the 1.5 μm laser transition to self-terminate. This is the reason why we grew and examined the decay time of 3 H4 and 3 F4 level of Tm 3+ /Pr 3+ :BGO crystals.

The lifetime of Tm 3+ : 3 H4 level almost keeps constant in Tm,Pr:BGO crystals with the Pr 3+ doping concentration from 0.1 to 0.3 at.%. Furthermore, when the Pr 3+ doping concentration is 0.5 at.%, the lifetime of Tm 3+ : 3 H4 level slightly dropes to 160 μs, which is 26.6% shorter compared with that of Tm:BGO crystal (218 μs). The comparable lifetime of Tm 3+ : 3 H4 level in Tm,Pr:BGO crystals suggests that co-doping Pr 3+ ions has little influence on the upper laser level 3 H4, On the contrary, the lifetime of Tm 3+ : 3 F4 level in Tm,Pr:BGO crystals decreases with increasing the doping concentration of praseodymium. The measured lifetime of Tm 3+ : 3 F4 level in 0.5 at.% Pr, 2 at.% BGO crystal is 360 μs, which is 93.6% shorter compared than that of Tm:BGO (5.64 ms), confirming that Pr 3+ ions can be used as efficient deactivation ions to depopulate the 3 F4 level of Tm 3+ for 1.5 μm emission.

The energy transfer efficiency (η) of Tm 3+ : 3 F4 level to Pr 3+ : 3 F2 level can be evaluated by the following formula [START_REF] Zhang | Enhanced 1.4 µm emissions of Tm 3+ via Tb 3+ deactivation in (Gd0.5Lu0.5)2SiO5 crystal[END_REF]:

η=1-τTm/Pr/τTm (1)
where τTm/Pr and τTm are the lifetime of Tm 3+ : 3 F4 level in Tm:BGO crystals with and without Pr 3+ co-doping, respectively. With increasing the Pr 3+ doping concentration from 0.1 at.% to 0.5 at.%, the energy transfer efficiency increased from 78.0% to 93.6%, which is much higher than that in Tm,Tb:GdScO3 (35.7% [START_REF] Li | Growth and spectroscopic properties of Tm 3+ and Tb 3+ co-doped GdScO3 crystal[END_REF]) and Tm,Tb:GLSO (74.4% [START_REF] Zhang | Enhanced 1.4 µm emissions of Tm 3+ via Tb 3+ deactivation in (Gd0.5Lu0.5)2SiO5 crystal[END_REF]), indicating that Pr 3+ ions can efficiently quench the lower laser level of the Tm 3+ : 3 F4. Moreover, the upper laser level ( 3 H4) to lower laser level ( 3 F4) fluorescence lifetime ratio in Tm,Pr:BGO crystals are calculated and are also listed in Table 1. The value of 0.5 at.% Pr, 2 at.% BGO crystal is 0.44, which is about eleven times that of the Tm:BGO crystal (0.04). All the results shows that the co-doping of Pr 3+ ions can lead to an accelerated depletion of population in the laser terminal level 3 F4, resulting in lower pump intensities to reach inversion.

Conclusion

Tm 3+ /Pr 3+ co-doped BGO crystals with various Pr 3+ doping concentration were successfully grown by the micro-pulling-down method. The BGO accepts Pr 3+ , Tm 3+ dopants which they do not make modifications of the stability and the local lattice structure. The efficiency of co-doping Pr 3+ to deactivate the the laser terminal level 3 F4 of the 3 H4 → 3 F4 transition of Tm 3+ in BGO crystal around 1.5 μm and thus to avoid the bottlenecking effect was demonstrated. It was also demonstrated that the Pr 3+ doping has little effect on the upper laser level 3 H4 of the 3 H4 → 3 F4 transition of Tm 3+ , which is beneficial to accelerated depletion of population in the laser terminal level 3 F4. Moreover, the energy transfer efficiency from Tm 3+ : 3 F4 level to Pr 3+ : 3 F2 level in 0.5 at.% Pr, 2 at.% Tm:BGO crystal was calculated to be 93.6% and the upper laser level ( 3 H4) to lower laser level ( 3 F4) fluorescence lifetime ratio (0.44) is about eleven times that of the Tm:BGO crystal (0.04). All of these results indicate that Tm,Pr:BGO crystal could be a promising laser material for 1.5 μm laser applications under the pump of a commercial 793 nm LD.

Funding

National Natural Science Foundation of China (No. 61621001) and "Qinglan Project" of the Young and Middle-aged Academic Leader of Jiangsu College and University.

Disclosures. The authors declare no conflicts of interest.

Fig. 1 .

 1 Fig. 1. Photographs of the as-grownTm,Pr codoped BGO crystals

Fig. 2 .

 2 Fig. 2. Room temperature absorption spectra of Tm:BGO and Tm,Pr:BGO crystals

Fig. 3 (

 3 Fig. 3 (a) NIR and (b) MIR fluorescence spectra of Tm:BGO and Tm,Pr:BGO crystals

Fig. 4 .

 4 Fig. 4. Simplified energy level diagram of Tm 3+ and Pr 3+ co-doped system.

Fig. 5 .

 5 Fig. 5. The decay curves of 3 H4 multiplet of Tm:BGO and Tm,Pr:BGO crystals

Table 1 .

 1 The fluorescence lifetimes of 3 H4 and 3 F4 levels in Tm:BGO and Tm,Pr:BGO crystals and the energy transfer efficiency from Tm 3+ to Pr 3+

	Crystals	τ( 3 H4)	τ( 3 F4)	η(ET)	τ( 3 H4)/τ( 3 F4)
	2 at.% Er:BGO	218 µs	5.64 ms	-	0.04
	0.1 at.% Pr, 2 at.% Tm:BGO	213 µs	1.24 ms	78.0%	0.17
	0.2 at.% Pr, 2 at.% Tm:BGO	194 µs	0.71 ms	87.4%	0.27
	0.3 at.% Pr, 2 at.% Tm:BGO	219 µs	0.79 ms	86.0%	0.28
	0.5 at.% Pr, 2 at.% Tm:BGO	160 µs	0.36 ms	93.6%	0.44


Fatima-Ezzahra Er-Rami

Marie Duffiet

Sean Hinkle

Jérémie Auvergniot

Maxime Blangero

Pierre-Etienne Cabelguen

Kyeongse Song

François Weill

Claude Delmas

Dany Carlier
email: dany.carlier@icmcb.cnrs.fr

Understanding the role of Al doping of LiCoO2 on the

Introduction

O3-LiCoO2 was first proposed as positive electrode material for lithium-ion batteries (LIBs) by Goodenough et al. in1980. 1 It is still widely used as positive electrode material in commercialized LIBs for portable electronics devices. Indeed, the main advantages are: a relatively large theoretical specific capacity as high as 274 mAh/g and an electrode density close to 4.0g/cm [START_REF] Orman | Cobalt(III) Lithium Oxide, CoLiO2: Structure Refinement by Powder Neutron Diffraction[END_REF] resulting in high practicable energy density and an excellent cycle stability. [START_REF] Wang | Reviving Lithium Cobalt Oxide-Based Lithium Secondary Batteries-toward a Higher Energy Density[END_REF] Layered LiCoO2 synthesized at high temperature exhibits the α-NaFeO2 type structure [START_REF] Orman | Cobalt(III) Lithium Oxide, CoLiO2: Structure Refinement by Powder Neutron Diffraction[END_REF] (R-3m space group with ahex.= 2.816 Å; chex. = 14.052 Å) with an O3 stacking according to the nomenclature developed by Delmas et al. [START_REF] Ohzuku | Solid-State Redox Reactions of LiCoO2 (R3m) for 4 Volt Secondary Lithium Cells[END_REF] Considerable efforts have been carried out to understand the structural changes occurring during electrochemical cycling by X-Ray diffraction and first principles calculations many years ago. [START_REF] Reimers | Electrochemical and in Situ X-Ray Diffraction Studies of Lithium Intercalation in LixCoO2[END_REF][START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF][START_REF] Van Der Ven | First-Principles Investigation of Phase Stability in LixCoO2[END_REF] When Li is deintercalated from LiCoO2, the material undergoes a series of phase transitions; the first occurs at a characteristic voltage plateau of ∼3.9V versus Li + /Li (0.75 ≤ x ≤ 0.94 in LixCoO2). It corresponds to a two-phase region induced by an insulator-metal transition where a newly appeared rhombohedral phase (called O32) coexists with the pristine LiCoO2 (O31). [START_REF] Van Der Ven | First-Principles Investigation of Phase Stability in LixCoO2[END_REF] More recently, the composition domain of the plateau has been revisited for a "real" stoichiometric LiCoO2: only when absolutely no structural defects are present, the two-phase domain starts quite at the onset of Li deintercalation. [START_REF] Ménétrier | Really" Stoichiometric LiCoO2[END_REF] Then, when ~ 0.5 Li is deintercalated, the system undergoes a transition from rhombohedral O3 to monoclinic O'3 [START_REF] Delmas | Structural Classification and Properties of the Layered Oxides[END_REF] due to a lithium/vacancy in-plane ordering. [START_REF] Reimers | Electrochemical and in Situ X-Ray Diffraction Studies of Lithium Intercalation in LixCoO2[END_REF][START_REF] Ohzuku | Solid-State Redox Reactions of LiCoO2 (R3m) for 4 Volt Secondary Lithium Cells[END_REF][START_REF] Shao-Horn | Probing Lithium and Vacancy Ordering in O3 Layered LixCoO2 (x≈0.5)[END_REF] Li deintercalation in LiCoO2 is highly reversible when the cell is cycled up to 4.2 V vs Li + /Li explaining its excellent cyclability when used as positive electrode material in LiBs with a specific capacity of 137 mAh/g vs Li + /Li. In order to achieve a higher capacity and increase the energy density, cycling voltage above 4.3 V vs Li + /Li is required. However, working at high voltage leads to a deterioration of the cycling performances, [START_REF] Yano | LiCoO2 Degradation Behavior in the High-Voltage Phase Transition Region and Improved Reversibility with Surface Coating[END_REF] caused by structural instabilities, electrolyte degradation and Co dissolution. [START_REF] Yazami | Mechanism of Electrochemical Performance Decay in LiCoO2 Aged at High Voltage[END_REF][START_REF] Xia | Phase Transitions and High-Voltage Electrochemical Behavior of LiCoO2 Thin Films Grown by Pulsed Laser Deposition[END_REF][START_REF] Takamatsu | First In Situ Observation of the LiCoO2 Electrode/Electrolyte Interface by Total-Reflection X-Ray Absorption Spectroscopy[END_REF][START_REF] Pender | Electrode Degradation in Lithium-Ion Batteries[END_REF][START_REF] Birkl | Degradation Diagnostics for Lithium Ion Cells[END_REF][START_REF] Amatucci | Cobalt Dissolution in LiCoO2-Based Non-Aqueous Rechargeable Batteries[END_REF] More specifically, for charge voltages exceeding 4.5 V, LiCoO2 exhibits detrimental phase transformations: O3 to H1-3 and H1-3 to O1 implying slab glidings. [START_REF] Van Der Ven | First-Principles Investigation of Phase Stability in LixCoO2[END_REF][START_REF] Xia | Phase Transitions and High-Voltage Electrochemical Behavior of LiCoO2 Thin Films Grown by Pulsed Laser Deposition[END_REF][START_REF] Chen | Methods to Obtain Excellent Capacity Retention in LiCoO2 Cycled to 4.5V[END_REF] Consequently, internal strain builds up, leading to crack formation and particle pulverization. [START_REF] Takamatsu | First In Situ Observation of the LiCoO2 Electrode/Electrolyte Interface by Total-Reflection X-Ray Absorption Spectroscopy[END_REF][START_REF] Jiang | Atomistic Mechanism of Cracking Degradation at Twin Boundary of LiCoO2[END_REF][START_REF] Jiang | Origins of Capacity and Voltage Fading of LiCoO2 upon High Voltage Cycling[END_REF] Several strategies were investigated to improve the cycling stability of LiCoO2. [START_REF] Hu | Enhanced High-Voltage Properties of LiCoO2 Coated with Li[Li0.2Mn0.6Ni0.2]O2[END_REF][START_REF] Gu | Enhanced Cycling Stability of High Voltage LiCoO2 by Surface Phosphorylation[END_REF][START_REF] Cho | Improvement of Structural Stability of LiCoO2 Cathode during Electrochemical Cycling by Sol-Gel Coating of SnO2[END_REF] Among these, metal substitutions of cobalt (Al, Mg, Ti…) have been demonstrated to be a promising and effective method to enhance the electrochemical performances of LiCoO2 at high voltage. [START_REF] Wang | Reviving Lithium Cobalt Oxide-Based Lithium Secondary Batteries-toward a Higher Energy Density[END_REF][START_REF] Zhang | Trace Doping of Multiple Elements Enables Stable Battery Cycling of LiCoO2 at 4.6 V[END_REF][START_REF] Liu | Approaching the Capacity Limit of Lithium Cobalt Oxide in Lithium Ion Batteries via Lanthanum and Aluminium Doping[END_REF][START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF][START_REF] Levasseur | Combined Effects of Ni and Li Doping on the Phase Transitions in LixCoO2 Electrochemical and 7 Li Nuclear Magnetic Resonance Studies[END_REF] In 1998, Ceder et al. [START_REF] Ceder | Identification of Cathode Materials for Lithium Batteries Guided by First-Principles Calculations[END_REF] identified aluminum as a potentially attractive dopant: they observed that substituting Co with Al leads to a higher Li cell average voltage but suffers from a lower capacity at the first charge up 4.4V vs Li + /Li. [START_REF] Ceder | Identification of Cathode Materials for Lithium Batteries Guided by First-Principles Calculations[END_REF][START_REF] Jang | LiAlyCo1-yO2 (R 3m) Intercalation Cathode for Rechargeable Lithium Batteries[END_REF] The similar Al 3+ and LS-Co 3+ ionic radii (0.535 Å vs. 0.545 Å) facilitate the substitution while maintaining the layered structure. It has been also reported that the substitution of Co by Al was effective on suppressing cobalt dissolution at 4.5 V versus Li + /Li. [START_REF] Myung | Effects of Al Doping on the Microstructure of LiCoO2 Cathode Materials[END_REF] In these previous works, the authors speculate that the dopant element is homogeneously distributed within the Co sites. Here, the challenge is to characterize both the Al repartition and the deintercalation mechanisms in an Al doped-LiCoO2 sample with a low dopant concentration at a level of industrial relevance (4%).

Previous studies have shown that [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li Magic Angle Spinning Nuclear Magnetic Resonance (MAS NMR) is a key tool to discuss the Li/M stoichiometry [START_REF] Levasseur | Evidence for Structural Defects in Non-Stoichiometric HT-LiCoO2: Electrochemical, Electronic Properties and 7 Li NMR Studies[END_REF][START_REF] Levasseur | Oxygen Vacancies and Intermediate Spin Trivalent Cobalt Ions in Lithium-Overstoichiometric LiCoO2[END_REF] whereas [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR can be used to characterize the Al/Co distribution homogeneity in the slab, even with a low doping level. [START_REF] Duffiet | Probing Al Distribution in LiCo0.96Al0.04O2 Materials Using 7 Li, 27 Al, and 59 Co MAS NMR Combined with Synchrotron X-Ray Diffraction[END_REF][START_REF] Han | Understanding the Role of Temperature and Cathode Composition on Interface and Bulk: Optimizing Aluminum Oxide Coatings for Li-Ion Cathodes[END_REF] Several 4% Al-doped LiCoO2 materials with controlled Li/(Co + Al) stoichiometries synthesized by various solid state routes have been recently studied in our group. [START_REF] Duffiet | Probing Al Distribution in LiCo0.96Al0.04O2 Materials Using 7 Li, 27 Al, and 59 Co MAS NMR Combined with Synchrotron X-Ray Diffraction[END_REF][START_REF] Duffiet | Investigation of Structural Failure Mechanisms of LiCoO2 at High Voltage and Material Optimization through Aluminum Doping[END_REF] Using 7 Li, and their doping homogeneity were discussed. The materials were compared with a LiCo0.96Al0.04O2 sample prepared by co-precipitation that also exhibits homogeneous Al/Co repartition, but lower particles size. For extensive cycling at high voltage, the sample prepared by the solid-state route showed, however, better stability, since larger particles size minimizes the electrolyte degradation's detrimental effect. Therefore, a sample prepared by a solid state route with homogeneous Al/Co distribution and a Li/M ratio equal to 1.00 has been selected for the present study. In a first part, we report the structural characterizations of the 4% Al-doped LiCoO2 discussing its homogeneity. Then, we aim to understand the role of Al-doping on the high voltage degradation mechanisms involved during cycling, combining techniques with different spatial and sensitivity scales such as XRD, [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li, 59 Co, and [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR and TEM.

Materials and methods:

4% Al-doped LiCoO2 sample was synthesized in Umicore R&D center by a solid state route from a mixture of Al2O3, Co3O4 and Li2CO3 as described in more details in our previous work. [START_REF] Duffiet | Probing Al Distribution in LiCo0.96Al0.04O2 Materials Using 7 Li, 27 Al, and 59 Co MAS NMR Combined with Synchrotron X-Ray Diffraction[END_REF] The de-intercalated LixCo0.96Al0.04O2 (x<1) samples were later prepared by electrochemical deintercalation in Li cells following the procedure disclosed in the next paragraph of this experimental section. A reference material, stoichiometric LiCoO2 (st-LiCoO2) with an exact Li/Co = 1.00 stoichiometry has been also synthesized by a solid state route. In order to avoid the formation a slightly overlithiated compound, as overlithiation impacts the cycling properties of LiCoO2, a blend nominal Li/Co=0.98 ratio was used, and the precursors were heat treated at 1050°C for 12h. In such thermodynamic synthesis conditions, the reaction products are expected to be st-LiCoO2 with traces of Co3O4 impurity as will be discussed in the following. Subsequently, grinding and annealing at 800 ° C for 5 hours was carried out in order to ensure good homogeneity of the materials and ensure representative and consistent analytical assessment. Final experimental (Li/Co) and (Li/(Al+Co)) ratios obtained by Induced Coupled Plasma (ICP) for st-LiCoO2 and LiCo0.96Al0.04O2 materials are gathered in Electron diffraction patterns as well as images of the samples were obtained using a JEOL 2100 transmission electron microscope working at an accelerating voltage of 200kV. Prior to their introduction in the microscope, samples were de-agglomerated in a mortar with addition of DMC, then a few drops of the resulting suspension were deposited on a copper grid with a formvar carbon film. No asymmetry of the peaks is seen by XRD suggesting a homogeneous Al doping. [START_REF] Duffiet | Probing Al Distribution in LiCo0.96Al0.04O2 Materials Using 7 Li, 27 Al, and 59 Co MAS NMR Combined with Synchrotron X-Ray Diffraction[END_REF] As Co is partially substituted by Al, the ahex. lattice parameter reflecting the M-M bond length slightly decreases from 2.816 (1) Å to 2.814 (1) (even though the difference is very small it goes with the tendency expected from ionic radii : r Al 3+ = 0.535 Å, r Co 3+ = 0.545 Å). The chex. lattice parameter related to the interlayer distance, increases more significantly from 14.052 (1) Å to 14.062 (1) Å in agreement with the literature. [START_REF] Jang | LiAlyCo1-yO2 (R 3m) Intercalation Cathode for Rechargeable Lithium Batteries[END_REF] LiCo0.96Al0.04O2 and st-LiCoO2 materials exhibit both a single [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR signal very close to 0 ppm (Figure S2a) that indicates that no paramagnetic ions are present in the Li environment, [START_REF] Ménétrier | Really" Stoichiometric LiCoO2[END_REF] in good agreement with the target formulas. Indeed, for Li/M = 1.00, only diamagnetic LS-Co 3+ ions with the t2g 6 eg 0 electronic configuration are expected in the layered phase. Note that the signal of LiCo0.96Al0.04O2 is slightly broader than the one of st-LiCoO2, as several Li environments versus Co/Al are expected in that material. However, the resolution of the [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR experiment is too low to distinguish clearly the different contributions. [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR were additionally used to characterize the Al/Co local environments (Figure S2b andS2c). As discussed in previous works, [START_REF] Duffiet | Probing Al Distribution in LiCo0.96Al0.04O2 Materials Using 7 Li, 27 Al, and 59 Co MAS NMR Combined with Synchrotron X-Ray Diffraction[END_REF][START_REF] Gaudin | III) Effect on 27 Al NMR Chemical Shifts in LiAlxCo1-xO2[END_REF] The deep characterization of the LiCo0.96Al0.04O2 sample (XRD combined to [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li, [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR) indicates that the sample is diamagnetic, therefore with no Li excess, and that the Al repartition in the Co layer is homogeneous.

Results and discussion

Electrochemical properties and general overview

LiCo0.96Al0.04O2 and st-LiCoO2 were used as positive electrode in Li cells that were cycled at C/20 rate. The resulting cycling curves clearly show that the substitution of Co by a small amount of Al (4%) has a profound effect on the electrochemical behavior (Figure 1 and2).

When st-LiCoO2 is delithiated, several phase transitions occur as discussed in the introduction and reported in Figure S4. During the 1 st cycling curve of LiCo0.96Al0.04O2 at C/20, a voltage plateau is observed in the low voltage region, but seems less extended that the one related to the Mott insulator-metal transition [START_REF] Ménétrier | The Insulator-Metal Transition upon Lithium Deintercalation from LiCoO2: Electronic Properties and 7 Li NMR Study[END_REF] observed for st-LiCoO2 (Figure 1). The cycling curve of LiCo0.96Al0.04O2 also does not feature the monoclinic transition around x = 0.5 [START_REF] Reimers | Electrochemical and in Situ X-Ray Diffraction Studies of Lithium Intercalation in LixCoO2[END_REF][START_REF] Ohzuku | Solid-State Redox Reactions of LiCoO2 (R3m) for 4 Volt Secondary Lithium Cells[END_REF][START_REF] Shao-Horn | Probing Lithium and Vacancy Ordering in O3 Layered LixCoO2 (x≈0.5)[END_REF] nor the O3/H1-3 voltage plateau on charge. Another striking difference between the two cycling curves is the large polarization observed at the beginning of the discharge for st-LiCoO2, not present for the doped material. Overall, the Al-doped sample shows a larger irreversible capacity than st-LiCoO2 as previously discussed in the literature. [START_REF] Myung | Effects of Al Doping on the Microstructure of LiCoO2 Cathode Materials[END_REF][START_REF] Levasseur | Oxygen Vacancies and Intermediate Spin Trivalent Cobalt Ions in Lithium-Overstoichiometric LiCoO2[END_REF][START_REF] Han | Understanding the Role of Temperature and Cathode Composition on Interface and Bulk: Optimizing Aluminum Oxide Coatings for Li-Ion Cathodes[END_REF][START_REF] Duffiet | Investigation of Structural Failure Mechanisms of LiCoO2 at High Voltage and Material Optimization through Aluminum Doping[END_REF] The suppression of Li/vacancy ordering for x = 0.5, as Co 3+ is partially substituted by Al 3+

suggests that the driving force and ΔG formation energy for Li-ordering is rather modest and easily perturbed by minor doping entropy increase. It was also the case for overlithiated LiCoO2, where even a very little excess of Li found in the Co layers hinders the Li-vacancy ordering at x = 0.5. [START_REF] Levasseur | Oxygen Vacancies and Intermediate Spin Trivalent Cobalt Ions in Lithium-Overstoichiometric LiCoO2[END_REF] The absence of this phase transition further confirms that Al is successfully homogeneously introduced into the LiCoO2 lattice.

If charged up to 4.5V vs Li + /Li, the capacity difference between the two materials is negligible vs. typical coincell method accuracy and precision (198 mAh/g for st-LiCoO2 versus 196 mAh/g for LiCo0.96Al0.04O2), but when charged up to 5V, the Al-doped phase shows a much lower capacity (247 mAh/g) than the undoped one (258 mAh/g) since less Li is deintercalated at 5V from the latter (Figure 1). In order to have a general overview of the mechanisms involved during charging, XRD and NMR study were performed on deintercalated samples prepared using a galvanostatic charge (C/20 rate) applied up to the desired Li content (x) followed by a relaxation step (Figure S5). The XRD patterns of the LixCo0.96Al0.04O2 and LixCoO2 materials are shown in Figure S6 and S7. Special attention was given to the (003) diffraction peak evolution that allows to directly follow the changes in the interslab distance, related to the lithium amount and stacking modifications. In the two systems, upon deintercalation this peak shifts towards lower angles upon deintercalation up to x=0.5, and then shifts back to higher angles for lower Li amount. This denotes an expansion of the structure along the chex. followed by a contraction (Table S2) which is consistent with the literature [START_REF] Reimers | Electrochemical and in Situ X-Ray Diffraction Studies of Lithium Intercalation in LixCoO2[END_REF] and commonly observed during the Li de-intercalation of layered LixMO2 materials. [START_REF] Van Der Ven | First-Principles Investigation of Phase Stability in LixCoO2[END_REF] The 7 Li, [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR spectra evolution versus the Li content are shown in Figure 3 and Figure S8.

While Li is in diamagnetic environments for the two pristine LiCo0.96Al0.04O2 and st-LiCoO2 samples, giving rise to signals close to 0 ppm, hyperfine interactions due to the oxidation of LS-Co 3+ (t2g 6 eg 0 ) to LS-Co 4+ (t2g 5 eg 0 ) ions are expected for the deintercalated samples.

The results observed for LixCoO2 samples are in good agreement with the literature. [START_REF] Ménétrier | The Insulator-Metal Transition upon Lithium Deintercalation from LiCoO2: Electronic Properties and 7 Li NMR Study[END_REF][START_REF] Geng | Monitoring the Evolution of Local Oxygen Environments during LiCoO 2 Charging via Ex Situ 17 O NMR[END_REF] As x decreases at the early stage of the charge for 0.95 < x < 1 (Figure S8), the intensity of the 0 ppm signal decreases, due to a strong hyperfine interaction between the single electron of the formed LS-Co 4+ (t2g 5 eg 0 ) ions and adjacent Li nuclei. Ménétrier et al. [START_REF] Ménétrier | The Insulator-Metal Transition upon Lithium Deintercalation from LiCoO2: Electronic Properties and 7 Li NMR Study[END_REF] interpreted this intensity decrease by the localized character of the single electron of Co 4+ , found in one of the t2g orbitals pointing directly towards adjacent lithium 2s orbital through the common edge of the CoO6 and LiO6 octahedra. Upon further lithium removal (Figure 3), a very weak new signal appears at 61 ppm for x = 0.95 and increases in intensity at the expense of the 0 ppm signal. This signal then shifts to higher ppm values until x=0.5. Ménétrier et al. [START_REF] Ménétrier | The Insulator-Metal Transition upon Lithium Deintercalation from LiCoO2: Electronic Properties and 7 Li NMR Study[END_REF] attributed this new signal to a Knight interaction due to the participation of the 2s orbital of lithium at the Fermi level, in good agreement with the metallic character of LixCoO2 phases for x < 0.74. [START_REF] Ménétrier | The Insulator-Metal Transition upon Lithium Deintercalation from LiCoO2: Electronic Properties and 7 Li NMR Study[END_REF] NMR is thus sensitive to the change in nature of the hyperfine interaction from localized electronic spins (Fermi contact) to delocalized ones (Knight shift). For x < 0.5, the Knight shifted signal is displaced back towards lower ppm values as also observed by Shimoda et al. [START_REF] Shimoda | In Situ NMR Observation of the Lithium Extraction/insertion from LiCoO2 Cathode[END_REF] The global evolution of the Li signal position versus x results from an interplay between the change in the local electronic structure (number of formal Co 4+ ions and localized versus delocalized electron spin character) and the change of local geometric structure (overall Li-O-Co distances or angles evolution).

Figure S8 shows that the observation of the [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR signals is rapidly loss while deintercalating Li + ions : the signals vanish for x = 0.80 and x = 0.95, respectively for [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co NMR. These signal loss could come from a strong Fermi contact interaction (localized state) and a stronger contribution of the s orbitals of Al and Co at the Fermi level compared to Li (delocalized state). For those nuclei, either a very broad signal not detected in our conditions or a very fast relaxation beyond the NMR timescale are to be expected. Therefore, no information from [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR regarding the local structure in samples prepared at high voltage could be obtained.

Detailed study in the high voltage region

In order to discuss the effect of the state of charge (Li content) and of the cutoff voltage, samples have been prepared following two routes as described in the experimental section: using a galvanostatic charge applied up to the desired Li content (x) followed by a relaxation step (Figure S5) and using a galvanostatic charge applied up to the desired voltage followed by a floating step until the current reached ~ 0 mA (Figure S9). The first route allows comparing samples prepared at the same state of charge, but prepared at different voltages that might induce different electrolyte degradation and the second route allows comparing phases at equilibrium for a given voltage but with different Li content. In addition, this second route leads to more aged samples since they spend more time at high voltage during the floating step. The XRD patterns of LixCo0.96Al0.04O2 and LixCoO2 are provided in Figure S6. In the LixCo0.96Al0.04O2, system, the x=0.5 and x=0.4 samples can be indexed using a single O3 type phase. However, for x=0.3, the co-existence of two (00l) lines indicates the presence of two phases that can be both indexed using the R-3m space, suggesting the presence of O3 and H1-3 phases.

Those two phases are still present for x = 0.2, but, for this composition, the XRD pattern shows much broader line, which could be explained by the existence of stacking faults within the layered oxide. [START_REF] Casas-Cabanas | Microstructural Characterisation of Battery Materials Using Powder Diffraction Data: DIFFaX, FAULTS and SH-FullProf Approaches[END_REF] In the case of st-LiCoO2, a pure O3 phase is observed for 0.5 < x ≤ 0.3 and three phases coexist for x=0.2 namely O3, H1-3 and another phase with an intermediate interslab space that can results from an intergrowth of O3 and H1-3 stackings. Note that such in intermediate phase was not observed during operando measurements and might be formed only during the relaxation step. [START_REF] Duffiet | Influence of the Initial Li/Co Ratio in LiCoO2 on the High-Voltage Phase-Transitions Mechanisms[END_REF] In the two systems, no signature of the formation of the O1 phase (P3m1 space group with ahex = 2.822 Å and chex = 4.29 Å) phase has been observed for compounds prepared up to x = 0.2.

Figure 4a

and 4b shows the [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li NMR spectra of the deintercalated samples from st-LiCoO2 and from LiCo0.96Al0.04O2. The spectra of the Li0.5CoO2 and Li0.5Co0.96Al0.04O2 phases are rather similar with a Knight shifted signal located at 94 ppm and 112 ppm respectively. For lower Li content, however, the spectra differ significantly. In the LixCoO2 system, the signal shifts to lower ppm values with a strong decrease in intensity and a broadening of the signal for x=0.2 as seen on Figure 4a. [START_REF] Geng | Monitoring the Evolution of Local Oxygen Environments during LiCoO 2 Charging via Ex Situ 17 O NMR[END_REF] This signal is assigned to Li in the O3-type environment, expected to be similar in the H1-3 stacking locally if we assume that a negligible amount of Li is present in the O1 interslab space. The broad signal for x = 0.2 may be due to the co-existence of 3 phases for this materials as seen by XRD, that probably exhibit slightly different compositions leading to different shifted signals. No rapid exchange is therefore observed at the NMR timescale between those domains.

In the LixCo0.96Al0.04O2 system, the main signal assigned to Li in O3 and H1-3 phases also shifts to lower ppm values, and for x < 0.3, new contributions appear in the [-20;50] ppm range with significant intensity (Figure 4b). Signals located in 10-50 ppm could be attributed to Li in spinel-type environment as a first approximation according to Godillot et al. [START_REF] Godillot | Promising Nanometric Spinel Cobalt Oxides for Electrochemical Energy Storage: Investigation of Li and H Environments by NMR[END_REF] , this point will be discussed in detail below. Additionally, signals located around 0 ppm can be assigned to diamagnetic lithium environments from the CEI (Cathode Electrolyte Interphases). Note that the signal intensity of the Li0.2Co0.96Al0.04O2 is larger than expected from the composition, indicating the participation of the electrolyte decomposition parasitic reaction in agreement with the presence of significant signals of Li in the CEI.

LixCo0.96Al0.04O2 and LixCoO2 samples prepared at the same voltages

The Li content of the materials prepared with different cutoff voltages (4.5V, 4.6V, 4.7V and 4.8V) is difficult to estimate due to the parasitic reaction of the electrolyte degradation, but are respectively lower than 0.32, 0.17, 0.10, 0.09 for LixCoO2 (Figure S9a) and 0.32, 0.24, 0.18, 0.15 for LixCo0.96Al0.04O2 (Figure S9b). The last two compositions exhibit therefore lower Li contents than samples of the previous section. One can observe that for floating voltages Vf ≥ 4.6 V, the cell voltage dropped significantly (VOCV as opposed to Vf) when the cell was disconnected. Indeed, the highest reached Vocv is only 4.6V, which remains higher than the voltage observed for the materials prepared without floating (Figure S.1). A really similar behavior was observed for the LixCoO2 samples prepared in the same way, but the 4.6V voltage remains stable after disconnecting the cell (VOCV ≈ Vf), whereas the voltage drops to 4.51V for the Al doped compound. The different materials were then characterized by XRD (Figure S10) and [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR (Figure 4c and4d).

All diffraction peaks for materials prepared at Vf =4.5 V can be indexed based on the O3 structure with rather similar Li content (Li0.24Co0.96Al0.04O2 and Li0.27CoO2 if one neglects the electrolyte degradation at this voltage). For the materials prepared at Vf =4.6V, two peaks are observed for LixCoO2 at 2θ=19.76° and 2θ=20.34°, which correspond respectively to the (006)H1-3 and (00l)int diffraction lines of the H1-3 and intermediate phase between H1-3 and O1 type structures as detected also by operando XRD. [START_REF] Duffiet | Influence of the Initial Li/Co Ratio in LiCoO2 on the High-Voltage Phase-Transitions Mechanisms[END_REF] Only the (006)H1-3 diffraction line is observed in the case of the Al-doped sample at the same Vf =4.6V. At 4.7V, three very broad contributions are observed for LixCoO2 assigned to the (006)H1-3, (00l)int and (001)O1 diffraction lines, indicating structures with a high density of stacking faults, resulting in overall structural disorder. In the case of the Al doped compounds, the (006)H1-3 peak remains narrow and the (001)O1 diffraction line appears at 2θ=20.92° for Vf ≥ 4.7V. Finally, LixCoO2 prepared at 4.8V also exhibits broad diffraction line assigned to a defected H1-3 phase with x ≤ 0.09 and LixCo0.96Al0.04O2 exhibits broad diffraction lines assigned to H1-3 and O1 phases with an overall Li content x ≤ 0.15.

From [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR (Figure 4c and4d), spectra of the materials prepared at Vf=4. The [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR results suggest that the CEI formed at high voltage could be either very thin or unstable for st-LiCoO2, and may be dissolved by the electrolyte solvent after the cell disconnection or by the electrode washing step using DMC prior to our measurements.

Additionally, no signal associated to a lithiated spinel phase was detected in this case. As opposed to st-LiCoO2, new local Li environments potentially assigned to both a disordered (as not seen by XRD) spinel-like lithiated phase and a stable CEI were observed for LixCo0.96Al0.04O2. These features could explain the higher stability upon cycling of the Al-doped samples and their highest irreversible capacity during the first cycle.

Structural characterization after cycling at high voltage (3 cycles)

To further understand the performance of LiCo0.96Al0.04O2 at high voltage, this material was charged-discharged 3 times between 3V and 4.8V with a C/20 rate then recovered at the end of discharge after a floating at 3V (Figure S11a). The sample was then characterized by XRD, MAS NMR and TEM. XRD and NMR data of this new sample was compared to the data collected for the pristine phase and a Li0.98Co0.96Al0.04O2 prepared by galvanostatic charge, as the phase at this composition is found to have cell parameters very close to the ones of the material recovered after 3 cycles (Figure S11b and Table S3). The XRD patterns of the 3 compounds, that can be indexed using a single phase in the R-3m space group of the O3 stacking. From the electrochemical curve (Figure S11a), a x=0.95 value was predicted, but, the cell undergoes some electrolyte decomposition at high voltage that shifted slightly the cycling curves on the left. TEM was used to confirm the hypothesis of the existence of a spinel phase in the cycled LixCo0.96Al0.04O2 material and to spatially localize it. Figure 7a, shows a nano-beam electron diffraction pattern obtained from the center of one crystal of the cycled LixCo0.96Al0.04O2 material (Figure 7c). This diffraction pattern corresponds to the [211] zone axis of the LiCo0.96Al0.04O2 phase using the hexagonal axis system (R-3m space group). When the nano-beam is displaced from the center to the edge of the crystal (Figure 7f), another pattern is obtained (Figure 7d). The presence of a spinel type phase has to be considered to explain it. Indeed, all the reflections of this pattern can be indexed by considering the [112] spinel zone axis as demonstrated by the superposition of the experimental pattern with the [112] spinel zone axis simulated pattern (Figure 7e). Nevertheless, note that some reflections have much higher intensities than others. These intense reflections seem to match those observed in the center of the crystal and indeed, some interplanar distances are identical in both structures as for instance d1-20lay and d4-40spi or d-204lay and d-1-11spi. The pattern obtained on the edge of the crystal must then be considered as the superposition of patterns belonging to both lamellar and spinel structures. This finding is thus in line with our [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR study indicating the presence of spinel-like environments for Li and additionally reveals that those domains are present at the surface of the particles only.

Discussion and conclusions

LiCo0.96Al0.04O2 was synthesized by a specific solid state route. [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR and XRD studies revealed that the Al-doping in this material is homogeneous while [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR indicates that its stoichiometry (Li/M=1.00) is ideal. Our electrochemical cycling study in Li cells reveals that even a low substituted amount of Co with Al (4%) has a profound effect on the electrochemical performance and yields a better cycling stability in the 3-4.6 V voltage range. XRD and [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR studies revealed that up to 4.5 V vs Li + /Li, the overall structural evolution of the LiCo0.96Al0.04O2 and st-LiCoO2 phases is similar, although i) the insulator-metal transition does not occur in the same composition range and ii) the monoclinic distortion at x = 0.5 is hindered by Al doping. At higher voltages (> 4.6 V vs. Li + /Li), noticeable structural differences have been evidenced in the deintercalated LixCo0.96Al0.04O2 and LixCoO2 materials and may be responsible or H1-3 structure were observed for LixCoO2, we believe that the CEI formed on that material has a different nature and is unstable (either dissolved after the cell disconnection or by the electrode washing step prior to the measurements). [START_REF] Zhang | Dynamic Evolution of Cathode Electrolyte Interphase (CEI) on High Voltage LiCoO2 Cathode and Its Interaction with Li Anode[END_REF] The different reactivity of the surface of LiCoO2 and Al-doped LiCoO2 has already been reported in literature: the basic character of the LiCo1-xAlxO2 surface makes these materials less reactive than LiCoO2 toward acidic species (such as HF) that are present in LiPF6-based electrolytes. [START_REF] Dahéron | Possible Explanation for the Efficiency of Al-Based Coatings on LiCoO2: Surface Properties of LiCo1-xAlxO2 Solid Solution[END_REF] The strong polarization observed at the beginning of the discharge from 5V in the st-LiCoO2 system might therefore result from i) the high resistivity character of that CEI at high voltage that may dissolve then while discharging as unstable; ii) the formation of the H1-3 or O1 phase at the particles surface with a low interslab space that may also prevent a good Li + ions diffusivity back to the particles.

Based on our study, we believe that Al-doping improves the cycling performance of LiCoO2 up to 4.6V, mainly by modifying the particles surface evolution upon cycling at high voltages.

Even during the first charge, it leads to the formation of a surface reconstruction with the formation of a lithiated spinel-type phase that should have a good ionic conductivity. In the case if LiCoO2, the formation of a spinel phase has been also reported, [START_REF] Wang | Electron Microscopic Characterization of Electrochemically Cycled LiCoO2 and Li(Al,Co)O2 Battery Cathodes[END_REF] is either not lithiated as not detected by [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR or formed only upon extended cycling. The lithiated spinel phase formed on the Al-doped compound, may then react differently with the electrolyte to form a stable CEI layer, whose composition is different from the one formed on LixCoO2. The formation of the Licontaining spinel phase and a stable CEI at the surface of the particles probably also explains the strong reduction of Co dissolution observed at high voltage as reported by Myung et al. [START_REF] Myung | Effects of Al Doping on the Microstructure of LiCoO2 Cathode Materials[END_REF]

ASSOCIATED CONTENT Supporting Information

XRD patterns and [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li, [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMR of the two pristine samples together with a comparison with statistical expected intensities for [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al NMR; First Cycling curve of st-LiCoO2 indicating the different phases; Cycling conditions used for Figure 2; Galvanostatic charge curves for the preparation of the deintercalated LixCoO2 and LixCo0.96Al0.04O2 phases; XRD patterns and cell parameters of the deintercalated LixCoO2 and LixCo0.96Al0.04O2 phases; [START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li, [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al and 59 Co MAS NMRS spectra of the LixCo0.96Al0.04O2 phases for 0.73 ≤ x ≤ 1; Galvanostatic charge curves followed by a floating step at different voltages for the preparation of the deintercalated LixCoO2 and LixCo0.96Al0.04O2 phases; XRD patterns of the phases prepared at a given voltage V ≥ 4.5V;

  A recycle time of D0 = 40s was used for st-LiCoO2 and LiCo0.96Al0.04O2 samples, whereas a shorter D0 = 4s was enough for all deintercalated samples, to avoid T1 saturation effects. Single pulse[START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR spectra were recorded on a Bruker 500 MHz spectrometer at 130.33 MHz (11.7 T magnet). The spectral width was set to 0.5 MHz and the recycle time to D0 = 5 s, long enough to avoid T1 saturation effects. As[START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al is a strong quadrupolar nucleus with I = 5/2, a short pulse length of 1.1 µs corresponding to a π/12 pulse determined using an aqueous 1 M Al(NO3)3 solution was employed. In these conditions, all of the -½ → +½ central transitions are equally excited regardless of the magnitude of the nuclear quadrupole coupling constants and one can extract quantitative data. As all samples contain only 4% of Al, overnight experiments (10240 scans) were carried out to ensure a good signal/noise ratio.59Co MAS NMR spectra were recorded on a Bruker 500 MHz spectrometer at 120.35 MHz (11.7 T magnet). The spectral width was set to 0.5 MHz and the recycle time to D0 = 1 s. A single pulse sequence using a short pulse length of 1.1 µs corresponding to a π/16 pulse was used to extract quantitative data. However, it requires a first-order phasing process with a sin x/x baseline correction due to the dead time of the spectrometer not easily determined due to large overlapping signals. The external reference was a 1M K3Co(CN)6 aqueous solution.

3. 1 .

 1 Figure S1. The two materials exhibit similar line-width and diffraction lines indexed in the rhombohedral R-3m space group of the O3 stacking. Note that the presence of Co3O4 as impurity in st-LiCoO2 is in good agreement with the Li/Co=0.98 ratio used for the synthesis and confirm the successful preparation of st-LiCoO2 (with no Li in the Co-layer, further confirmed by NMR below).

Figure 1 :

 1 Figure 1: First charge-discharge curves in the 3-5 V voltage range vs. Li + /Li of LiCo0.96Al0.04O2 and st-LiCoO2 (C/20 cycling rate).

Figure 2 :

 2 Figure 2 : Discharge profiles (3V-4.6V) of Li//LiMO2 (M=Co, Al) cells observed for the 7, 8, 31 and 32 th cycles of a) st-LiCoO2 and c) LiCo0.96Al0.04O2 using cycling conditions provided in Table S1. Corresponding discharge capacities versus cycle number of b) st-LiCoO2 and d) LiCo0.96Al0.04O2.

Figure 3 :

 3 Figure 3: Ex-situ 7 Li MAS NMR spectra for a) LixCoO2 and b) LixCo0.96Al0.04O2 recorded using a rotor-synchronized Hahn echo sequence (vR = 30 kHz). Arbitrary intensities are used here to better visualize the shifts.

3. 3 . 1

 31 LixCo0.96Al0.04O2 and LixCoO2 samples prepared at the same states of charge a) Li x CoO 2 b) Li x Co 0.96 Al 0.04 O 2

Figure 4 :

 4 Figure 4: Comparison of 7 Li MAS NMR for highly deintercalated phases of a) LixCoO2 and b) LixCo0.96Al0.04O2 with fixed Li contents of c) LixCoO2 and d) LixCo0.96Al0.04O2 with fixed floating voltages. The ex situ NMR spectra intensities were normalized with respect to the mass of the sample analyzed and number of scans.

  5V are very similar to the ones of, respectively, Li0.3CoO2 and Li0.3(Co0.96Al0.04)O2 phases, with a single Knight shifted signal located around ~75 ppm. For the three samples prepared at larger floating voltages Vf ≥ 4.6V, the line shape of the 7 Li signal of LixCoO2 samples remains almost unchanged (Figure 4c). Several contributions are observed for LixCo0.96Al0.04O2 samples (Figure 4d) as already discussed above, but in different relative amounts. The intensity of the signal located in the [~50-130] ppm range expected for Li in the O3 and H1-3 phases strongly decreases while the intensity of signals corresponding to Li in (supposedly) spinel type environment at [10-50] ppm and to Li in the CEI around 0 ppm remain essentially the same.

Figure 5 :Figure 6

 56 Figure 5: a) Ex-situ 7 Li solid-sate MAS NMR of the LixCo0.96Al0.04O2 and b) zoom of the spectra.

Figure 7 :

 7 Figure 7: Nanobeam electron diffraction on the cycled LixCo0.96Al0.04O2. a) Experimental diffraction pattern related to the center of the particle indexed with the LixCo0.96Al0.04O2 structure. b) Comparison of experimental and simulated patterns. c) Location of the beam corresponding to pattern a. d) Experimental diffraction pattern related to the edge of the particle

  for the enhanced cycling stability of Al-doped LiCoO2. The XRD study shows that the deintercalated phases prepared for x < 0.2 from st-LiCoO2, either by fixing the Li content or by a floating method, present the co-existence of several phases (O3, H1-3, O1 and phases with intermediate interslab space) with stacking faults. Whereas, deintercalated samples prepared from LiCo0.96Al0.04O2 are more crystalline, and the formation of phases with intermediate interslab space is not observed. In addition, at 4.6V, the O1 phase is not observed for the Al-doped sample.More interestingly,[START_REF] Ven | First-Principles Evidence for Stage Ordering in LixCoO2[END_REF] Li MAS NMR reveals very different Li local environments in the high voltage phases. Unlike LixCoO2, for Al-doped phases, signals assigned to Li in spinel-like environments and in a diamagnetic CEI layer are observed in addition to the signal corresponding to Li in the interslab space of the O3 /H1-3 structures. Those observations do not depend on the materials preparation (by fixing the Li content or the floating voltage). The spinel phase is still detected after a few cycles in discharged LiCo0.96Al0.04O2 phase is and the TEM study revealed that it is located at the surface of the particles. Since only signals of Li in the interslab space of the O3

Table

.

  The results show that the experimental values are in good agreements with the theoretical ones.

	Target compounds homogeneous samples stable at high voltage. All sample powders of the deintercalated compounds Li/Co Li/M (M=Co+Al)
	Target were then split in two parts for NMR and XRD characterization. Experimental	Target	Experimental
	St-LiCoO2 (+Co3O4) Powder X-ray diffraction (PXRD) patterns were collected on a PANalitycal X'pert PRO 0.98 0.971 --
	Li1.00Co0.96Al0.04O2 MPD, diffractometer in Bragg-Brentano θ-θ geometry equipped with a Fe filter, a spinner and 1.04 1.039 1.00 0.998
	X'Celerator multi-strip detector using the Co-Kα. Radiations. Each pattern was recorded in the
	(10 -120°) 2θ range using 0.016° intervals and 15 h total counting time. XRD patterns of the Table 1: ICP analysis of the LiMO2 samples prepared using Li/Co = 0.98 and Li/M = 1.00 with M = 96% Co + 4% Al. deintercalated materials were collected on a PANalitycal X'pert 3 Powder diffractometer using the
	Cu-Kα. radiations, that allows the study of powder packed in 0.3 mm diameter capillaries.
	Electrodes of undoped and 4%Al-doped LiCoO2 were prepared by mixing 90% weight of
	NMR experiments were performed using a standard Bruker 2.5 mm MAS probe with a active material, 5% of carbon black, and 5% of polyvinylidene fluoride (PVDF). A slurry was formed by adding N-methyl-pyrrolidone (NMP) and mixed for 1h. The slurry casted onto a 30 µm-30 kHz typical spinning frequency. 7 Li MAS NMR spectra were recorded on a Bruker 300 Advance
	thick-aluminum foil, was dried in an oven for 2 hours at T = 80 °C, calendared and punched into spectrometer at 116.66 MHz (7.05 T magnet). A Hahn echo sequence [tπ/2-τ1-tπ-τ2] synchronized
	discs with an electrode density of ~3.0g/cm 3 . The electrodes were then dried overnight under with one period of rotor rotation was used for a 30 kHz spinning frequency. The 90° pulse duration
	vacuum at T = 120 °C and stored in an argon-filled glovebox, where cell assembling was carried was equal to tπ/2 = 2.0 µs and determined using a LiCl 1 M solution.
	out. Electrochemical tests were both carried out in coin cells and Swagelok type cells. Coin cells
	were used to assess the electrochemical performance of all samples (ie, charge and discharge
	capacities, cycle life performance, etc…), while the Swagelok-type design was favored when the
	cycled samples needed to be recovered for further characterization. Two different loading of the
	electrodes were used depending on the cell setup. The coin cells were assembled using
	LiCo0.96Al0.04O2 or st-LiCoO2 positive electrodes with a 10 mg/cm 2 active material loading, a disc
	of Viledon® and Celgard 2400® as separators and a Li metal sheet as negative electrode with
	1 M LiPF6 in EC: DEC: DMC as electrolyte. A higher loading of ~ 30 mg/cm 2 was employed to
	prepare Li de-intercalated LixCo0.96Al0.04O2 or LixCoO2 materials in the Swagelok setup to have a
	sufficient amount of powder for all ex-situ characterizations. One more Viledon separator was
	added in this case. The materials were recovered in an argon-filled dry box, washed with dimethyl
	carbonate (DMC) to remove residual electrolyte salt and dried under vacuum. Deintercalated
	LixCo0.96Al0.04O2 and LixCoO2 samples were prepared by two different routes. Route 1: a
	galvanostatic charge at a C/20 rate applied up to the desired Li content (x) followed by a relaxation
	step; Route 2: a galvanostatic charge at a C/20 rate applied up to the desired voltage followed by a
	floating step until the current reached ~ 0 mA. This second procedure was used to prepare

  at ~62 ppm corresponds to Al surrounded by 6 Co in the slab (denoted as Al-(Al0Co6)). This signal exhibits a clear second order quadrupolar line shape. The other signals located at 55 ppm, and 48 ppm correspond to Al surrounded by n Al 3+ and (6-n) Co 3+ with n=1 and 2, respectively. Signals for Al in a Al-rich environment with n ≥ 3, expected at lower ppm value[START_REF] Levasseur | Evidence for Structural Defects in Non-Stoichiometric HT-LiCoO2: Electrochemical, Electronic Properties and 7 Li NMR Studies[END_REF] were not detected in our sample due to the very low intensity predicted for those environments if Al distribution is homogeneous. Indeed, the relative intensities observed in our sample are really close from the one expected for a statistical repartition of Co and Al in the slab (See FigureS3).St-LiCoO2 exhibits a single signal located at 14137 ppm as expected for LS-Co 3+ surrounded by 6 LS-Co 3+ in the slab, in agreement with literature[START_REF] Siegel | Co and 6,7 Li MAS NMR in Polytypes O2 and O3 of LiCoO2[END_REF] . Despite broader signals observed for59 Co MAS NMR, compared to[START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR, distinct signals can be resolved for LiCo0.96Al0.04O2. The most intense one located at 14132 ppm is assigned to LS-Co 3+ ions surrounded by six other LS-Co 3+ ,

	59 Co MAS NMR spectra of LiCo0.96Al0.04O2 and st-LiCoO2 materials are displayed in Figure S2c.

a set of signals is observed, each corresponding to a different Al environment denoted as Al-(AlnCo6-n) for Al surrounded by n Al 3+ and (6-n) Co 3+ . The signal with the higher chemical shift located denoted as Co-(Co6Al0). The other signals located at 14105 ppm, 14077 ppm, 14047 ppm, and 14021 ppm are assigned to Co surrounded by n Al 3+ and (6-n) Co 3+ in the slab with n varying respectively from 1 to 4, denoted as Co-(Co6-nAln). 34 Comparison with statistical intensities was less straightforward due to the strong overlap of those broad signals, but no Al-rich environments (n = 5 and n = 6) were detected.

Al , 59 Co MAS NMR and synchrotron X-Ray Diffraction (XRD) the exact Li/M stoichiometry

Three first cycles (3V-4.8V) of LiCo0.96Al0.04O2 and XRD pattern of the LixCo0.96Al0.04O2 phase recovered at the end of discharge compared to the pristine and the deintercalated Li0.98Co0.96Al0.04O2.

for highly deintercalated materials and were attributed to Li in spinel-like environment as a first approximation. According to Godillot et al. [START_REF] Godillot | Promising Nanometric Spinel Cobalt Oxides for Electrochemical Energy Storage: Investigation of Li and H Environments by NMR[END_REF] the 12 ppm and 24 ppm signals are assigned to Li in the 8a tetrahedral sites of the spinel structure, that may differ from their Co/Al local environments here and the 43 ppm signal is assigned to Li in the 16c octahedral sites. No secondary phase was detected by XRD indicating that the spinel-type phase is either very disordered phase or only present at a very local scale (Figure S11b). 59 Co and [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR is not really informative regarding the spinel formation since a loss of signals for those phases is expected : in Co3O4, an electronic delocalization occurs on Co ions leading to a very strong hyperfine interaction with 59 Co nuclei that directly carries the electronic spins; in CoAl2O4, the hyperfine interaction between the Co 2+ ions and the [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al nuclei is also very strong and the absence of signal for Al was already reported [START_REF] Gaudon | Synthesis Temperature Effect on the Structural Features and Optical Absorption of Zn1-xCoxAl2O4 Oxides[END_REF] and was also confirmed by our group while heat treating Co3O4 with Al2O3 at 1000 °C. [START_REF] Gaudin | III) Effect on 27 Al NMR Chemical Shifts in LiAlxCo1-xO2[END_REF] Indeed, the 59 Co and [START_REF] Zhang | The Origin of Heavy Elemental Doping to Relieve the Lattice Thermal Vibration of Layered Materials for High Energy Density Li Ion Cathodes[END_REF] Al MAS NMR spectra of the LixCo0.96Al0.04O2 sample recovered after 3 cycles up to 4.8V are really similar to the ones of the x= 0.98 sample, so no major structural modification occurs within the O3 phase (Figure 6).

AUTHOR INFORMATION

ACKNOLEDGEMENTS

The authors wish to thank Umicore company for financial support and fruitful discussions and Cathy Denage for experimental support and SEM image acquisition.


Dominique Planson
email: dominique.planson@insa-lyon.fr

Camille Sonneville
email: bcamille.sonneville@insa-lyon.fr

Pascal Bevilacqua
email: cpascal.bevilacqua@insa-lyon.fr

Luong Viêt

V Phung

Besar Asllani
email: ebesar.asllani@insa-lyon.fr

Dominique Tournier

Pierre Brosselard

4H-SiC PiN diode protected by narrow field rings investigated by the micro-OBIC method

Keywords: 4H-SiC bipolar diode, high voltage device, electro-physical characterisation, periphery protection, OBIC

Introduction

Silicon Carbide (SiC) is an attractive semiconductor material for high power and high temperature applications. To reach the expected breakdown voltage, an effective junction termination is required to spread the electrical field profile towards edges mitigating the crowding effect. Among different techniques used for periphery protection [1], Junction Termination Extension (JTE) is a well-known technique often chosen due to their relative design simplicity. However, the structure of the present paper relies exclusively on guard rings, a more challenging approach where small technological deviations on the rings position can completely negate the benefits of the periphery protection. To get the most from those rings, key parameters such as the distance between each ring and the number of rings must be carefully chosen. The efficiency of this periphery protection is optimal when the electric field at the edge of each ring is nearly equal for all rings. Optical Beam Induced Current (OBIC) technique is a non-destructive characterization technique used here to analyze the electric field distribution at the edge of the device periphery. OBIC is an alternative technique compared to other complicated methods like Kelvin Probe Force Microscopy (KPFM) and Electronic Beam Induced Current (EBIC) microscopy, for high resolution analysis of electronic devices. This technique has been used to investigate silicon diodes protected by field rings [START_REF] Stengl | High-voltage planar junctions investigated by the OBIC method[END_REF], with spacing greater than 10 µm, as well as 4H-SiC PiN diodes protected by JTE [3], [4]. In this paper, micro-OBIC technique is applied to SiC high voltage bipolar diodes where the design of the guard rings is more tightly constrained to analyze the experimental behavior of the periphery protection with respect to the applied reverse voltage.

Experimental Setup

In the micro-OBIC experimental setup, a 349 nm UV pulsed laser is used to generate electron-hole pairs (EHPs) in the sample due to the laser beam with energy greater than the band gap energy. A set of properly adjusted optics (mirrors, microscope objective …) is inserted on the beam path to finally get a focused laser spot with a diameter of about 1-4 µm as shown on Fig. 1. The sample is placed on a motorized stage and the position is controlled with LabView (i.e. on X, Y, Z automatically and inclination teta manually). The laser beam energy at the top surface of the device is high enough to generate electron-hole pairs (EHPs), so that an OBIC current could be measured. The micro-OBIC principle has already been described in [4].

The testbench allows us to realize a spatial mapping of OBIC signal, either on a (X, Y) plane or only on X (or Y) lines. For X or Y lines, the step between two displacements of the laser beam is 200 nm. The DUT can be reverse biased by applying voltage up to 500V thanks to the SMU (Source Measure Unit) Keithley 237. The cross section of the DUT and its periphery is sketched on the Figure 2. For the sake of clarity, only one periphery protection is sketched in Fig. 2-a). The epilayer is 14.8 µm thick, n-type doped (6´10 15 cm -3 ). The main junction is realized by ion implantation, at the same time as the field rings, which were carefully designed by Caly Technologies. The diode is then protected by a passivation layer on the front side.

The measurements were realized on two different square-shaped bipolar diodes protected by field rings, namely device 1 and device 2, with and without passivation layer respectively. The metallized anode contact is square-shaped, with a side length of 200 µm x 300 µm, and a corner curvature radius of 200 µm as shown in Fig. 2-b.

This non-destructive technique allows to measure OBIC currents by scanning the outer edge of the biased diode, in order to assess the periphery protection efficiency. OBIC current increases as the electric field intensity increases in the structure. There is no OBIC current on the metallization, neither far from the periphery of the diode. These OBIC current profiles would come in handy to the designer to optimize the breakdown voltage. The diode is protected by 22 field rings as shown in Fig. 2, with increasing spacing between them (0.07 µm from one ring to the next). The distance between the main junction and the first ring is 1.2 µm. All the rings have the same width (1.5 µm).

Inside a vacuum probing chamber, on-wafer measurements of the devices demonstrated a stable blocking voltage up to 2 kV as shown in Fig. 3.

Results and Analysis

By applying different reverse voltages (up to 500V) to the same diode and re-scanning the very same location, several OBIC currents can be compared to each other, as shown in Fig. 4. As the reverse voltage increases, the OBIC current increases for the same location, due to the increase of the electric field. One can also observe the extension of the space charge region as the reverse voltage increases and the role of outer field rings, witness of the peripheral protection good design. As the reverse voltage increases, the number of OBIC peaks increases. From 20 µm up to 97 µm, each maximum is related to the presence of a ring as described in [1] and, more precisely, is located on the outer edge of each ring junction. As the reverse voltage increases, the presence of the rings is more clearly revealed by µ-OBIC. Indeed, the number visible rings by OBIC increase with the reverse voltage. These measurements up to 500V, show the presence of very close rings (between 1.2 µm for the first two rings and 2.67 µm for the last two rings). One can count all the 22 field rings for a reverse voltage of 500V plus one more peak for the border of the main junction. That means that at 500V all the rings are revealed by the OBIC method. The amplitude of the maximum doesn't increase much from one voltage to another. The distance between the peaks is in total accordance with the layout of the rings as shown in the figure and in Table 1. Table 1 gives the position of the first 11 rings based on its left lateral edge that is oriented towards the main junction when viewed from the cross-section depicted in fig. 2a. The OBIC signal amplitude is given for each left lateral edge. Comparing the data from Table 1 and the OBIC profile shown in Fig. 4, each dip in the profile perfectly matches with the left lateral edge of each ring. All the data in the table 1, were extracted from diode #2 biased under 500 V. After measuring diode #1, some scratches appeared at the surface of the device under test, due to the sweeping of the laser. In fact, this device was passivated with a polyimide layer, allowing a uniform and homogeneous spreading of the electric field as shown in Fig. 4, and the opportunity to test the DUT without package, i.e. avoiding any electrical arcing in the air. With respect to the beam penetration, this additional layer acts as a UV light filter and then can be physically altered, consequently the OBIC signal is reduced by a factor of 100 approximately. Nonetheless, Figure 4 shows the benefits of polyimide layer at the top in terms of electric field distribution: at 500V the OBIC current is almost constant whereas Figure 5 shows a decrease in OBIC current value. For diode #1 one can observe all the rings from 300V.

Fig. 4. OBIC currents at the periphery of the diode device #1 (with polyimide) for different reverse voltages (ranging from 0V up to 500V). The mask layout with the spacing between the rings is in grey color.

In order to study the periphery protection without any passivation layer, the polyimide was removed with a very high-density plasma on a new die. Then, the same experiment has been performed on diode #2 without polyimide layer at the top. The applied reverse voltage was also limited to 500V. The results are shown in Fig. 5. For diode #2 One can observe all the rings from 400V.

Compared to the previous measurements on both devices, performing measurements without the polyimide layer gives OBIC profile that reflects the peripheral protection more regular. In both cases, the peaks are related to the distance between the rings, and that is true even for very small distance as low as 1.2 µm. More investigation is needed to properly discriminate the electrical effect of the passivation layer and its optical effect as a light filter.

Fig. 5. OBIC currents at the periphery of the diode device #2 (without polyimide) for different reverse voltages (ranging from 0V up to 500V by 50V step). The mask layout with the spacing between the rings is in grey color.

Fig. 1 .

 1 Fig. 1. Schematic representation of the micro-OBIC testbench.

Fig. 2 .

 2 Fig. 2. a) Schematic cross section view of half vertical PiN diodes protected by 22 field rings (principle is sketched here only with 13 rings), b) photograph of the diode with the 50 µm (diameter) gold ball bonding

Fig. 3

 3 Fig. 3 Reverse characteristics of both devices (diode #1 with polyimide passivation, diode #2 without any passivation)

Table 1 :

 1 Distance from the main junction for ring mask layout and respective OBIC signal values.

	Ring	#1	#2	#3	#4	#5	#6	#7	#8	#9	#10 #11
	number
	Distance	1.2	3.97	6.81	9.72	12.7	15.75 18.87	22.06	25.32 28.65 32.05
	from main

Summary

This paper presents OBIC measurements performed on field rings protected bipolar diodes at different voltages. This test-bench has a well improved spatial resolution compared to the previous OBIC test-bench. We estimated theoretically and experimentally the spatial resolution of the micro-OBIC to be around 1-4 μm. A better resolution is interesting to characterize peripheral protections, such as JTE rings, whose lateral extension is of the order of magnitude of the micron. JTE rings are directly observable in these experiments, even with narrow distances between each 22 rings. In the near future, this study will be completed by measurements at higher reverse bias under vacuum and to be closed to the breakdown voltage of the device (1 700 V). Indeed, we could transfer the µ-OBIC test bench into a vacuum 10 kV chamber. It would be also interesting to study these same diodes by other complementary characterization methods such as micro-Raman spectroscopy, SEM/EDX and TOF-SIMS.