H Ichou

M Rossi

T Belgrand

Assessments of High-Power Solid State Transformers based on Grain-Oriented magnetic cores

Keywords: Grain Oriented Electrical Steel, Medium Frequency Transformer, Solid-State Transformer (SST)

All fields of activity using electric energy are more than ever challenged for efficiency and versatility of energy flow at reasonable costs. The emergence of electronic components with high voltage and current capabilities enables to cope with those challenges. The paper deals with assessments on Medium Frequency (MF) high-power Solid State Transformer (SST). A good technical-economic balance can be achieved by assembling suitable high power SST cells made with mature technologies for power electronics, magnetic cores and simple and reliable control strategies. In this framework, the study and design of a SST based on elementary cells involving a Grain Oriented Electrical Steel (GOES) core is addressed.

I. INTRODUCTION

S OLID State transformers (SST) are getting more and more popular due to increased ability of semi-conductor components to switch high voltages and currents as well as in the development of strategies for controlling the reversible power flow in electric grids. This growing technology, already used for HVDC [START_REF] Jianyang | Hvdc bidirectional power flow using solid state transformer[END_REF], is promising for smart grids development [START_REF] Yao | Mu synthesized robust controller for multi-sst islanded smart grid[END_REF] but also for shrinking down the size of the components of grid nodes. The grid stability need nodes able to control the energy flow, which makes key assets the SSTs. It may also be used to reduce the volumes of converters for e-mobility, railway traction, substations or marine applications [START_REF] Zhao | Power electronic traction transformer-medium voltage prototype[END_REF] [START_REF] Dujic | Power electronic traction transformer technology[END_REF].

The transformer is a critical element to the SST-cell, in that it enables the voltage change as well as the galvanic separation between input and output converters. Many research works have been performed on SST involving transformers based on amorphous or nanocrystalline magnetic cores [START_REF] Zhao | A practical solution of high-frequencylink bidirectional solid-state transformer based on advanced components in hybrid microgrid[END_REF] [START_REF] Du | Design considerations of high voltage and high frequency transformer for solid state transformer application[END_REF]. In these latter cases, power converters are requiring high level technologies and complex topologies [START_REF] Zhao | A practical solution of high-frequencylink bidirectional solid-state transformer based on advanced components in hybrid microgrid[END_REF] [START_REF] Du | Design considerations of high voltage and high frequency transformer for solid state transformer application[END_REF].

Small size SST means high frequency and high flux density level in magnetic cores. Multilevel topologies are of growing interest [START_REF] Iturriz | Phase-controlled multilevel converters based on dual structure associations[END_REF] [8] [START_REF] Kjaer | A primary-switched line-side converter using zero-voltage switching[END_REF], because they are more prone to reproduce the classical sinusoidal case and are build to avoid strong dV /dt's adverse effects. Conversely the use of two or three levels rectangular wave-forms is far more simple for highpower applications. A methodology for analyzing the overall losses of medium frequency transformers in a SST cell supplied by non-sinusoidal voltages have been studied in [START_REF] Villar | Global loss evaluation methods for nonsinusoidally fed medium-frequency power transformers[END_REF]. Several approach of losses estimation can be used, such as the Improved Generalized Steinmetz Equation (IGSE) [START_REF] Aguglia | Laminated magnetic materials losses analysis under non-sinusoidal flux waveforms in power electronics systems[END_REF].

The paper analyses the possibilities for designing highpower SST cells made with well experienced components. The transformer is built around a wound core made of thin Grain Oriented Electrical Steel (GOES) sheets adapted to medium frequencies (MF) and high peak flux density. GOES are mature products, available on the market and offering widths up to 1m, high saturation levels (2.03T) and high temperature capabilities (up to 400 • C without major loss of its magnetizing ability) thanks to its mineral insulation layers. Shaped into a wound core, the electrical steel is then used such as the magnetization process occurs parallel to the easy Rolling Direction magnetization axis; avoiding any extra losses due to corners of a stacked structure.

The proposed assessments consider both the SST cell structure and the transformer properties. The cooling system design is out of the scope of the paper.

II. SOLID STATE TRANSFORMERS CELL

A. SST cell structure

The basic concept of the cell is shown in Fig. 1. The SST cell consists of a Medium-Frequency Transformer (MFT) and two single phase converters. The power transfer P of a MFT depends on its volume, on frequency f , on flux density B m and on thermal balance [START_REF] Multon | Composants passifs (magnétiques et capacitifs) de l'électronique de puissance[END_REF]. The maximum power flow in the MFT is given by [START_REF] Jianyang | Hvdc bidirectional power flow using solid state transformer[END_REF], where the core cross-section A e and the winding area A w corresponds to the MFT volume [START_REF] Multon | Composants passifs (magnétiques et capacitifs) de l'électronique de puissance[END_REF]. K B is window filling factor, K i is the crest factor, and J the current density in the conductor, which depends on the global thermal balance.

P = K i .K B .J.A e .A w .f.B m (1)
From ( 1), the size of the transformer can be expressed as a function of frequency and peak flux density :

A e .A w = P f.B m .K i .K B .J (2)
The decrease of the MFT size is obtained by increasing the frequency or/and by increasing the peak magnetic induction level. This has a strong influence on specific loss [START_REF] Basso | Power losses in magnetic laminations with hysteresis: Finite element modeling and experimental validation[END_REF] and consequently a good balance between all the parameters has to be found.

When combined together to transfer higher power values, the SST cells are connected in series to the High Voltage (HV) grid and in parallel on Low voltage (LV) one. The converters connected in series are placed in side insulated racks, control signals are transferred via optic fiber that ensure commonmode insulation. The MFT primary winding common-mode insulation level must be reinforced compared to the secondary one connected to the LV grid. Fig. 2 shows the MFT topology.

B. Estimation of losses in the converters

The details of the SST cell are presented in Fig. 3 where each IGBT leg operates only one switching per period imposing rectangular voltages v 1 and v 2 . The voltage and current waveforms are shown in Fig. 4. The time shift of v 2 relatively to v 1 controls the power flow. Fig. 4 is drawn for a time shift that corresponds to the maximum power flow from grid 1 toward grid 2.

The MFT voltage ratio is set to 1 and the two dc grid voltage levels are supposed to have their amplitudes equal to each other (v A = v B ). For the core size defined in Fig. 2, the magnetizing current is much lower than the load current, corresponding to the energy transfer from a grid to the other, because of the very high magnetic permeability of the wound GOES core.The peak magnetizing current in less than 0.5A for every configurations considered in section E that correspond to peak load current in the range of 80A -100A.

As the magnetizing current can be neglected in a first approach, the current waveform depends on the instantaneous voltage difference v 1 -v 2 and on the leakage inductance L K of the MFT when the resistances are neglected.

L K di LK dt = v 1 -v 2 (3)
The second graph of Fig. 4 shows the voltage difference (v 1v 2 ) and the resulting current i LK flowing across the windings considering the proposed SST cell parameter settings. The corresponding IGBT bridge input current i dc1 is drawn in the third part of Fig. 4.

The average input current < i dc1 > is calculated by (4) from the DC input voltage V A and the rated SST cell power. The maximum value in the input I M in the inverter and the transformer is derived form the current wave-forms in figure 4 by [START_REF] Zhao | A practical solution of high-frequencylink bidirectional solid-state transformer based on advanced components in hybrid microgrid[END_REF].

< i dc1 > = P out v A (4)
I M = 2. < i dc1 > (5)
In the electronic converters, two kinds of losses must be considered: switching and conduction ones. The conduction losses are given by ( 6) from the saturation voltage V CE(sat) and the collector current mean value < i c > calculated from the current wave-forms presented in the lower part of Fig. 4.

P cond = V CE(sat) . < i c > (6)
< i c > = 1 T T 0 i c dt = 5 16 I M (7)
The losses in the converters are calculated separately for the waveforms shown in Figure 4.

The energy W OF F lost during the transistor turn-off transition is given by [START_REF] Oliveira | A two-stage ac/dc sst based on modular multilevel converterfeasible to ac railway systems[END_REF].

W OF F = 1 2 .v A .I M .t f (8)
The same estimation can be made for the IGBTs of the second converter. The total switching losses of SST cell can be estimated by [START_REF] Kjaer | A primary-switched line-side converter using zero-voltage switching[END_REF]. Where t f is the fall time.

P SW = W OF F .f = 1 2 .v A .I M .t f .f (9)
The two converters are identical to each other, then, the respective losses are equal. Since each converter is made of four identical switches, total converter loss of SST cell is given by [START_REF] Villar | Global loss evaluation methods for nonsinusoidally fed medium-frequency power transformers[END_REF]. The data-sheet of IGBTs module gives

t f = 60ns, V CE(sat) = 2.1V and diode with V F = 1.8V at 125 • C [14]. P conv = 8.(P SW + P cond,IGBT + P cond,diode ) (10)

C. Winding losses estimation

For medium and high frequencies, skin and proximity effects naturally increase the wire AC resistance relatively to the DC one. Several work show that the AC resistance increase is limited by using litz wires made of many elementary enamelled wires of small diameter connected in parallel. The ratio R AC /R DC depends on the number of elementary enamelled wires, on their diameter on the winding topology and on the frequency [START_REF] Sullivan | Cost-constrained selection of strand diameter and number in a litz-wire transformer winding[END_REF] [START_REF] Wojda | Winding resistance of litz-wire and multi-strand inductors[END_REF]. The manufacturers propose several products of round or rectangular litz wires designed for getting good filling factors [START_REF] Vonroll | Winding wires and litz wires[END_REF]. The product choice depends on the technological choices made for the MFT design. For getting a reasonable degree of freedom for litz wire choice in the manufacturer data-sheets, the maximum ratio of R AC /R DC has been set to 1.4. The DC resistance values (R DC p and R DC s ) are computed form lengths and cross section of the wire used for manufacturing the primary and secondary windings. These resistances depend on the sizing options presented in section E. The RMS currents in the MFT windings are computed from the wave-form of the middle graph of Fig. 4 applying the classical definition [START_REF] Aguglia | Laminated magnetic materials losses analysis under non-sinusoidal flux waveforms in power electronics systems[END_REF] and the copper losses derive from this current by [START_REF] Multon | Composants passifs (magnétiques et capacitifs) de l'électronique de puissance[END_REF].

I RM S = 1 T T 0 i 2 LK dt = I M 2 3 ( 11
)
P W = (R ACp + R ACs )I 2 RM S (12)

D. Core losses

The manufacturing process of the core on the magnetic properties is of utmost importance as well as the influence of the voltage waveform. This is why two kinds of voltage waveform have been assessed, the classical sine and the rectangular one. Experimental investigations have been made at several frequencies and peak flux densities for determining the core losses. Fig. 5 show a wound core, which has been cut for facilitating the transformer coil assembly. For testing, two specific winding have been distributed along the core. The numbers of turns are identical (N 2 = N 1 ). The primary winding is fed by a square wave or sine wave. The steps to get the core loss are:

• Measure the i LK (t) current flowing through primary winding when the secondary winding in open circuit. This current is measured with a current probe, it is in direct related with the magnetic field according to the Ampere's law i LK (t) = H(t).lm N1 .

• Measure the voltage induced in the secondary winding

v 2 (t); v 2 (t) = -A c .N 2 dB(t)
dt . • Calculate the average of the product i LK (t).v 2 (t) during a period T given by ( 13):

P = 1 T N 1 N 2 T 0 i LK (t).v 2 (t) dt (13)
Measurements are performed at steady state after demagnetizing the core. The primary current is measured using a non-inductive shunt with a channel of the oscilloscope and the voltage with another one. Data are stored in an oscilloscope, 10000 points for each measurement, and transferred to a PC; and a Matlab script perform the integration.

For the same peak flux density, measurement results of Fig. 6, made at 1 and 2kHz, show lower losses for square voltages than for sine ones. Similar result exists in the literature on ferrite and amorphous cores [START_REF] Chen | Comparisons of high frequency magnetic core losses under two different driving conditions: A sinusoidal voltage and a square-wave voltage[END_REF]. These results show that the square voltage waves obtained with simple converters using only a single switching per IGBT leg and per period can be used rather than to more complex PWM inverters able to produce voltage waves close than sine ones. Fig. 7 shows core losses for square-voltage waves versus frequency for several peak flux densities.

E. SST cell global design

From sections B, C and D, the overall efficiency of SST cell, η, can be calculated by (14) for several design options.

η = P out P out + P conv + P W + P core (14)
The design procedure of the MFT consist of several steps with two constants for getting an easy comparison: • the wound-core size presented in figure 2 (core weight = 2.9 kg).

• the copper filling ratio set to 30% for getting room for the insulation(that corresponds to a copper weight = 10.18 kg.

The other variable parameters are tuned for analyzing the design balance when the cooling system is able get temperatures compatible with the temperature electrical insulation system. Four input variables are used :

• the rated power;

• the input SST-cell voltage • the frequency;

• the maximum flux density. The lines of Table I are computed from the input constants and variables with the results of sections B, C and D. The turn number N of each coil is computed from the maximum flux density, the core section and the frequency. The cross section (S cu ) of the Litz copper wire depends on the size of the core window because the filling factor is kept constant for insulation. Consequently the current density and the copper losses depends on the input variables. The converter losses is slightly dependent on frequency and more on the current. The core losses depends strongly on the frequency. As the filling factor is kept constant, the wire section increase with frequency and the the winding losses reduce. This approach provides a global analysis on the balance of losses for three frequencies.

Table I, shows that an overall efficiency over 95% can be achieved when the cooling system keeps the temperature under the electrical insulation thermal class. Table II estimates the SST-cell performances in such cases with a higher voltage for 3kHz and two peak flux densities.

Comparing to the results of Table I, Table II show that a higher power/mass ratio is possible accepting slightly higher copper losses. For the same rated power a higher voltage reduces slightly the losses in the converter. Efficiency remains around 96% whatever the balance between core losses and copper losses. These results show that it is possible to design high power SST-cells with a power/mass ratio up to 4.6kW/kg with wound cores made of thin GOES able to operate at high temperatures.

The next step is to define the cooling system and a thermal design able to guarantee a winding temperature compatible with the thermal class of present-day polymers that can reach 200 • C for a live span of 20000 hours.

III. DETAILS ON CONTROL STRATEGIES

A. Model of a cell within the global structure

The main concept of a modular SST based upon the association of N elementary cells is illustrated by Fig. 8. The input converters are connected in series to share the high voltage (HV) and low voltage side connects the converters in parallel for providing a high current. The control of each cell ensure the global balance of the structure.

Neglecting the magnetizing current, the transformer can be modelled by the equivalent circuit (Fig. 9) the reduces to the figure Fig. 10 for the equivalent currents and voltages seen from the primary side. R S is the equivalent global resistance (R s = n 2 .r 1 + r 2 ) and L K the global leakage inductance

(L K = n 2 .l 1 + l 2 ). R s .i LK + L K . di LK dt = v 1 -v 2 ( 15
)
Fig. 8: Modular structure of a high power SST made with N identical cells.

Fig. 9: Equivalent circuit of the transformer of a cell.

For handling the control of the IGBTs, the principle described by [START_REF] Byen | A high-efficiency variable modulation strategy for a dual-active-bridge converter with a wide operating range[END_REF] has been chosen. It consists in monitoring the three control parameters defined in Table III for the case of one commutation per period as showed by Fig. 11. Each converter generates a positive, a negative and a zero voltages. The operating principle depends on the phase shift φ introduced between the rectangular voltages generated by the two bridges. If the two voltages v 1 and v 2 are exactly the same, there will be no power transfer. However, as soon as a phase shift is introduced between them, an alternating current is generated. This current will depend on the voltage difference, as well as on the leakage inductance L K of the transformer.

B. Different types of control and comparison 1) Phase Shift Modulation: (PSM)

The power transfer is controlled via the single parameter φ, with:

β 1 = β 2 = 0 and -π 2 ≤ φ ≤ π 2 .
It is inversely proportional to the leakage inductance as expressed in ( 16). This equation is representative of the magnitude and direction of the power flow by the value of φ. The maximum transmitted power is obtained for φ= ± π 2 .

P = V A .V B ω.L K .π φ.(π-| φ |) (16)
2) Trapezoidal Modulation: (TZM)

Parameters β 1 , β 2 can be used to as complementary for controlling the power flow control adjustment as illustrated by Fig. 11. In this graph, the three parameters vary, for computing the transmitted power. The maximum power is found when φ = π 2 and β 1 ,β 2 = 0; which corresponds to the PSM method.

3) Triangular Modulation: (TRM)

This method is a particular case of TZM. Only β 1 , β 2 are varied, keeping φ = 0. Several statements arise:

• β 1 = β 2 : The two control v 1 and v 2 voltages are identical, no current in the AC link thus no power. • β 2 ≥β 1 : Power flows from primary to the secondary.

• β 2 ≤β 1 : Power flows from the secondary to the primary.

The maximum power transmitted for this method is half the one obtained with the other ones.

The phase shift modulation (PSM) and the trapezoidal modulation (TZM) show a better control strategy compared to triangular modulation (TRM), with regard to the maximum transmitted power. However, for a complex association of several SST cells (Fig. 8), the global control is complex and the other parameters β 1 and β 2 provide degrees of flexibility that may be useful.

C. Choice of a control strategy for the elementary cell

The simpler control strategy, that uses only one parameter φ, has been chosen for the SST cell. This PSM method is adequate to control an SST cell with a wide power transfer range. Fig. 12 presents a practical application where the SST cell supplies a DC load, which must be regulated at a constant voltage v B whatever the load current i Load . The output voltage v B is kept constant by the SST cell that acts on its internal control parameter. The control loop of Fig. 13 gives details. The controller acts on the internal parameter φ for regulating the output voltage v B very near the set point value whatever could be the SST cell output current and the input voltage in the operating range of the cell. The output voltage v B depends on the output capacitance and currents. Assuming that the capacitance C 2 has a large value, v B is suppose to remain constant during short times. Consequently the instantaneous value for the control loop can be assimilated to the mean value of the voltage defined at the short time scale ot the power electronic converters (v B ≈ < v B >). The output power is expressed by [START_REF] Chen | Comparisons of high frequency magnetic core losses under two different driving conditions: A sinusoidal voltage and a square-wave voltage[END_REF] and the output current mean value by [START_REF] Byen | A high-efficiency variable modulation strategy for a dual-active-bridge converter with a wide operating range[END_REF].

v B = 1 C 2 (< i dc2 > -i Load )dt (17)
P 2 =< v B > . < i dc2 > (18) < i dc2 >= V A ω.L K .π φ.(π -φ) (19)
The controller of Fig. 13 must be tuned for providing the the loop stability and the required voltage accuracy of the SST cell.

IV. DISCUSSIONS ON THE RELIABILITY OF THE PROPOSES SST CELLS

The main goal of the paper consist in opening the possibility of designing large power elementary SSTs cells, in the range of 1M W , for building large power SSTs by assembling these elementary cells. For such high power systems, the structural reliability is a key point. The option is to base the design on mature technologies (TRL8 or TRL9) [START_REF] Herr | Lopak igbt modules-inverter reliability benefits[END_REF].

The structural reliability of the SST cell depends on the one of the subsystems (Inverter and transformer) and on the quality of their assembly:

• The inverters operate in kHz range frequency with Silicon IGBTs and simple commands (two switching per period). • The medium frequency transformer designed with a hot wound core made of thin GOES laminations and coils made of standard Class 200 • C enamelled wire. • The mechanical and electrical assembly of components uses mature methods used for a long time in power electronic industry.

A. Inverter structural reliability

Generally speaking, the lifetime of Si-IGBT modules depends mainly on electrical and thermal stresses. The reliability relative to the electrical stress is obtained, using a correct safety controlled with a correct safety margin for the module rated voltage and low inductance connections (laminated busbars and low parasitic inductance decoupling capacitors placed at the strategic position of the laminated busbar).

The Si-IGBTs structural reliability relative to the thermal stress has been studied for a long time [START_REF] Onuki | Reliability of thick al wire bonds in igbt modules for traction motor drives[END_REF]. It is now well known that IGBT module can work for limited lifetimes over the rated chip temperature [START_REF] Morris | Operating igbts above rated junction temperature limits: Impacts to reliability and electrical performance[END_REF]; with the proposed simple command, the IGBT losses are limited and the cooling system can be designed for limiting the IGBT chips under the rated temperature for increasing the expected lifespan.

The structural reliability of IGBT modules depends mainly of the mechanical stress on the bonding and on the packaging due to thermal cycles [START_REF] Choi | Power cycling test and failure analysis of molded intelligent power igbt module under different temperature swing durations[END_REF]. This study is made with short thermal cycles for pratical experiments reasons, but it can be extended to long cycles. Figure 14 extracted from [START_REF] Choi | Power cycling test and failure analysis of molded intelligent power igbt module under different temperature swing durations[END_REF] gives an example of results for 10s thermal cycles and a end-of life criterian define as an increase of the Collector-Emitter saturation voltage. This study shows that the sample lifespan is over 10000 cycles.

B. MFT structural reliability

The important parameter is the thermal class of the winding wire; it is defined by the IEC 60216-3 standard [START_REF]Détermination des propriétés d'endurance thermique des matériaux isolants électriques[END_REF] as the maximum temperature for a lifetime of 20000h with a failure rate of 50 % for an applied voltage of 1000V during 1min. This standard uses Arrhenius' law for predicting the lifetime of a coil. The IEC 60172 standard [START_REF]Méthode d'essai pour la détermination de l'indice de température des fils de bobinage émaillés[END_REF] The hot GOES wound core is also an element of the SST cell structural reliability. The natural insulation layer of GOES products is made of an inorganic magnesium silicate film (glass film) on which has been deposited a second film composed of inorganic components based on phosphates and silicates [START_REF] Lemaître | Matériaux magnétiques doux cristallins. acier électrique à grains orientés[END_REF]. This Insulation configuration is said EC-5 on EC-2 according to the IEC 60404-1-1 standard [START_REF]Magnetic materials-part 1-1: Classification-surface insulations of electrical steel sheet, strip and laminations[END_REF]. GOES withstands sheet annealing at 850 • C. This layer is extremely stable at the expected operating temperatures of the SST cell.

MFT cooling is a key point of its structural reliability. The GOES wound core can operate at high temperature for getting a good power/mass ratio, but the winding temperature must remain under its thermal class. Consequently, a thermal insulation must be placed between the core and the winding.

V. CONCLUSION

The proposed assessments on high-power SST cells based on wound cores made of thin sheets of high performance grainoriented electrical steel show that high power/mass ratio can be reach with a correct global efficiency at medium frequencies and high flux densities. The characterization measurements show that the magnetic core losses in square voltage excitation are lower than in sine one for the same maximum flux density and frequency. Such voltage waveforms require simple converters that can be easily transposed to very high power SST cells.

The lifetime of high power IGBTs and winding wires is a necessary element to increase confidence in the design of SST cell high power. The next step consists to define a specific thermal management of the medium frequency transformer that guarantees a winding temperature compatible with its insulation system class with a wound core that is not practically limited in temperature.

ACKNOWLEDGMENT

This work is funded by ThyssenKrupp Electrical Steel.

Fig. 1 :

 1 Fig. 1: Topology of a SST cell based on a simple MFT.

Fig. 2 :

 2 Fig.2: Core and windings of the MFT with distances able to withstand 30kV for the primary coil common-mode insulation.

Fig. 3 :

 3 Fig. 3: SST structure of an elementary cell.

Fig. 4 :

 4 Fig. 4: Voltage and current wave-forms in the SST cell for a time lag of T 4 that corresponds to the maximum power flow.

Fig. 5 :

 5 Fig. 5: Cut wound core.

Fig. 6 :

 6 Fig. 6: Comparisons of magnetic core losses under sinusoidal and square voltages.

Fig. 7 :

 7 Fig. 7: Core losses as a function of frequency for several flux density peak values.

Fig. 10 :

 10 Fig. 10: Global equivalent of the SST cell circuit.

Fig. 11 :

 11 Fig. 11: Voltage and current waveforms, depending on the control parameters φ , β1 and β2.

Fig. 12 :

 12 Fig. 12: The SST cell in its functional context.

Fig. 13 :

 13 Fig. 13: The block diagram of a SST cell.

Fig. 14 :

 14 Fig. 14: Example of thermal cycling tests of IGBT modules. The endof life criteria is the unacceptable increase of the on voltage [23].

  allows the Arrhenius curve to be drawn for enamelled wire of class 200 • C often used for electric machine windings, a lifetime of 200000 hours can be achieved by working at a lower temperature (155 • C) according the Arrhenius law (figure 15).

Fig. 15 :

 15 Fig. 15: Arrhenius curve for thermal class 200 • C winding wire [25].

TABLE I

 I DESIGN PARAMETERS OF THE SST CELL AT 1.6T

		P = 40kW -800V	P = 50kW -800V
	f (kHz)	1.5	2	3	1.5	2	3
	N (turns)	128	96	64	128	96	64
	Scu(mm 2 )	17.5	23	35	17.5	23	35
	I RM S (A)	81.6	81.6	81.6	102	102	102
	Jcu(A/mm 2 )	4.65	3.49	2.35	5.8	4.3	2.9
	Pconv(W )	644	653	672	804	816	840
	Pcore(W )	371	596	1163	371	593	1163
	P W (W )	847	476	211	1321 744	330
	η	95.55 95.87 95.1	95.2	95.8 95.5


Andres Martinez

Vittorio Grimaldi

Christian De Vita

Kumar Deepak

Isabelle Sharma

Aurore Gallet

Juan Andrieux

Kamal Arocas

Laurent Hammani

Jean Claude Markey

Weeber

Deepak Kumar

Isabelle Gallet

Aurore Andrieux

Juan Arocas

Kamal Hammani

Laurent Markey

Jean-Claude Weeber

Alexandre Bouhelier

Marco Sampietro

Andrea Melloni

Francesco Morichetti
email: francesco.morichetti@polimi.it

D K Sharma

Inline Photothermal Surface Plasmon Detector Integrated in Titanium Dioxide Waveguides

come

Abstract-Densely integrated photonic integrated circuits (PICs) require efficient solutions for monitoring the light intensity on chip in order to implement control and configuration operations to set and stabilize the working point of the circuit. To this end, waveguides supporting the propagation of surface plasmon polaritons (SPPs) are good candidates to realize small-footprint light detectors. In this work, we report on the realization of an inline Surface Plasmon Detector (SPD) that exploits the photothermal effect to monitor the optical power in a titanium dioxide (TiO2) optical waveguide. Detailed design guidelines are provided to maximize the responsivity of the SPD, taking into account the effects of the metal geometry on the coupling between the dielectric and plasmonic modes, the power dissipated in the metal, and the equivalent thermal resistance of the structure. Experimental validation of the proposed device is provided demonstrating an ultra-compact 1.6-µm-long SPD operating at a wavelength of 1550 nm with a sensitivity of -20 dBm and a bandwidth higher than 100 kHz. The proposed device concept can be ported to generic dielectric platforms and to other wavelength ranges where SPP propagation is supported.

Index Terms-integrated photodetectors, integrated photonics, optical sensors, photothermal effects, surface plasmon polaritons.

I. INTRODUCTION

P hotonic Integrated Circuits (PICs) are evolving towards architectures with an ever-increasing scale of integration that implement more and more complex functionalities. To develop PICs that can be reconfigured reliably, on-chip photodetectors must be integrated to locally monitor the optical power and closed-loop control systems have to be designed to set and stabilize their working point. Semiconductor photodetectors are widely employed in the near-infrared range because of their small size, fast detection speed, high detection efficiency [START_REF] Hui | Photodetectors[END_REF], and monolithic integration on silicon and III-V photonic platforms. This family of transducers includes photodiodes [START_REF] Logan | Monitoring and Tuning Micro-Ring Properties Using[END_REF], photoconductors [START_REF] Kingstone | Optical Detectors[END_REF], sub-band-gap detectors based on defect state absorption [START_REF] Logan | Defect-Enhanced Silicon-on-Insulator Waveguide Resonant Photodetector With High Sensitivity at 1.55 μm[END_REF], and surface state absorption [START_REF] Morichetti | Non-Invasive On-Chip Light Observation by Contactless Waveguide Conductivity Monitoring[END_REF]. However, these approaches cannot be directly deployed onto dielectric waveguide platforms, such as silicon (oxy) nitride or other glass-based technologies.

A promising approach to implementing light sensing in small-footprint devices is offered by surface plasmon polaritons (SPPs), which are guided waves that can propagate along the surface of a conductor [START_REF] Homola | Surface Plasmon Resonance Based Sensors[END_REF] [START_REF] Barnes | Surface plasmon subwavelength optics[END_REF]. SPP waves remain tightly confined on the metal-dielectric interface because the free electrons of the conductor respond collectively by oscillating in resonance with the electromagnetic wave [START_REF] Barnes | Surface plasmon subwavelength optics[END_REF] [START_REF] Mayer | Localized Surface Plasmon Resonance Sensors[END_REF]. SPPs are only possible for wavelengths longer than a particular critical wavelength, which depends on the plasma frequency and is specific to each metal [START_REF] Homola | Surface Plasmon Resonance Based Sensors[END_REF] [7] [START_REF] Mayer | Localized Surface Plasmon Resonance Sensors[END_REF]. Moreover, SPPs are only TM polarized as there is no solution to the Maxwell equations that would represent a bounded mode for the TE polarization [START_REF] Homola | Surface Plasmon Resonance Based Sensors[END_REF]. As SPP propagation is associated with the absorption of optical power in the metal, a temperature-dependent change of the metal resistivity is generated by the photothermal effect, which can be exploited to monitor the intensity of the light in the waveguide [START_REF] Kumar | Power monitoring in dielectric-loaded surface plasmon-polariton waveguides[END_REF] [10] [START_REF] Gosciniak | Power monitoring in dielectric-loaded plasmonic waveguides with internal Wheatstone bridges[END_REF] [12] [START_REF] Weeber | Thermoelectric detection of waveguided surface plasmon propagation[END_REF]. Different configurations have been proposed for photothermal detectors such as a Wheatstone bridge [START_REF] Gosciniak | Power monitoring in dielectric-loaded plasmonic waveguides with internal Wheatstone bridges[END_REF] and hybrid plasmonic waveguides [START_REF] Wu | Ultracompact on-chip photothermal power monitor based on silicon hybrid plasmonic waveguides[END_REF].

In this work, we demonstrate an in-line Surface Plasmon Detector (SPD) integrated into a TiO2 platform that exploits the photothermal effect to monitor the optical power inside the waveguides. The device simply consists of a thin strip of gold placed underneath the TiO2 core of the waveguide, whose lightdependent resistance is measured by using a four-probe lock-in detection scheme. The paper is organized as follows. In Sec. II we illustrate the working principle of the device and the design rules to optimize the device performance in terms of insertion loss and photothermal responsivity. To this end, systematic numerical simulations based on a multi-physics 3D finite element method (FEM) including optical, thermal, and electrical analysis were carried out to investigate the impact of the metal strip geometry in the coupling between the dielectric TiO2 waveguide and the SPP waveguide, and to maximize the absorption efficiency in the metal. In Sec. III the technology employed to integrate the SPD on a TiO2 platform is described. Experimental results reported in Sec. IV demonstrate the functionality of the device both in terms of responsivity (-20 dBm) and speed (> 100 kHz). A concluding section summarizes the main achievement of this work.

II. DEVICE CONCEPT AND DESIGN OPTIMIZATION

The waveguide used for this work is realized with an aircladded 1.2 µm-wide and 370 nm-high titanium dioxide (TiO2) core on top of a 2 µm-thick silicon dioxide (SiO2) on a Si substrate. In the SPD region, a gold strip is inserted underneath the waveguide core, as shown in Fig. 1(a) and (b). At the transition between the dielectric waveguide (no metal strip) and the plasmonic waveguide, the TM fundamental mode of the dielectric waveguide may excite several SPP modes (see Fig. 2) [START_REF] Homola | Surface Plasmon Resonance Based Sensors[END_REF] [14] [START_REF] Mirzanejhad | Numerical study of nanoscale biosensor based on surface plasmon polariton propagation in Mach-Zehnder interferometer structure[END_REF]. SPP modes propagate in the plasmonic waveguide with different phase constants and accumulate different phase shifts for the same travelled distance. At the end of the metal strip, the relative phase between the SPPs modes modifies the shape of the overall electromagnetic field and hence the mode matching with the fundamental TM mode of the output dielectric waveguide. This constructive (maximal transmission) or destructive (minimal transmission) interference results in an insertion loss that depends on the metal strip length L. In the next sections, 3D FEM Multiphysics simulations are reported to analyze and optimize the coupling between the dielectric and the plasmonic modes. Factors such as the power dissipated inside the metal and the thermal behavior of the whole structure are investigated to maximize the responsivity of the device while keeping a low insertion loss for in-line detection of the guided light. All the reported simulations are carried out at a wavelength of 1550 nm.

A. OPTICAL ANALYSIS

The design of the in-line SPD detector has to be optimized in order to maximize the optical power absorption in the metal strip (this being at the origin of the photothermal effect), while keeping low the overall loss, meaning that other sources of loss (e.g., out of plane scattering and back reflections) must be minimized. To this end, full vectorial EM simulations were carried out to evaluate the dependence of the complex refractive index of the SPP modes on the transversal geometry of the metal strip. The refractive indices of the SiO2 bottom oxide and the TiO2 layer are 1.445 and 2.2, respectively, as confirmed by material characterization (see Sec. III). The cross-sectional shape of the fundamental mode of the SPP waveguide (TMpolarized) is shown in Fig. 3(a). The waveguide supports the propagation of higher order modes, as the one shown in Fig. 3(b), whose effects will be studied in the following. (1) which depends on the field intensity adjacent to the Au strip, increases for thinner strips. For thick strips, the evanescent component of the field penetrates more into the metal [START_REF] Okda | Sensitivity enhancement of a difference interferometer refractive index sensor based on a silicon-on-insulator hybrid plasmonic waveguide[END_REF], but the rapidly decaying SPP mode profile makes this contribution too small to affect the propagation loss α. Furthermore, if the metal is too thick, the significant discontinuity at the edges of the metal strip would severely increase back reflections and reduce the light coupling to the detector. Thus, for light monitoring applications, the metal thickness should be maintained thin enough to increase the gold absorptance while also minimizing back reflection loss. In the following, a thickness T = 34 nm is assumed. Also, we observe a decrease in both the real and the imaginary (absolute value) part of the refractive index versus the thickness of the gold. This behavior is typical of a short-range surface plasmon polariton [START_REF] Homola | Surface Plasmon Resonance Based Sensors[END_REF]. The large index contrast between the dielectric (2.2 for the TiO2 and 1.445 for the SiO2) introduces a larger cutoff frequency for the long-range modes, which in the considered structure are not supported [START_REF] Berini | Plasmon-polariton waves guided by thin lossy metal films of finite width: Bound modes of structures[END_REF].

Concerning the cross-sectional width W, the gold strip should be as narrow as possible to maximize the SPD responsivity. In this way, the equivalent thermal resistance of the detector increases, inducing a larger variation of the temperature and consequently of the read-out voltage (see Sec. B) for the same amount of dissipated power. Since the lateral confinement of the SPP modes is given by the width of the dielectric core (that in the considered case is 1.2 µm), there is a minimum width Wmin of the metal strip for which the guided light does not interact with the free electrons of the gold strip outside the core. As shown in Fig. 4(b), this implies that real (solid blue line) and imaginary (solid red line) parts of the effective index neff of the SPP mode remain constant for W > Wmin. Choosing a strip width slightly wider than Wmin guarantees the highest responsivity while keeping the detector robust to fabrication tolerances. For the considered structure, Wmin = 2.6 µm, so W was set to 3 µm. The length of the metal strip L is the fundamental design parameter as it determines the intensity of the transmitted optical wave and the amount of optical power that is absorbed by the metal. Fig. 5 shows the simulated normalized transmittance Pout across the SPP detector versus the length L (W = 3 µm and T = 34 nm). Results show that the TM transmission has a decaying sinusoidal behavior with a period of 2.4 μm. The exponential decay is due to the fact that the SPD has a fundamental mode with a complex-refractive index neff = 2.084 -i0.016 which corresponds to a propagation loss of 0.56 dB/µm. The periodic oscillation suggests that another SPP mode is excited, which propagates independently along the top and bottom surfaces of the metal at a different speed. Numerical simulations show that the waveguide supports the propagation of three SPP modes. In Fig. 3 (2)

In the considered SPP waveguide geometry (W = 3 µm and T = 34 nm), where ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,0 } = 2.084 and ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,2 } = 1.489, a beat length Δ𝐿𝐿 = 2.6 µm is found, which is in very good agreement with the result shown in Fig. 5. The minima in the transmitted power of Fig. 5 occur when the two interfering SPP modes are out of phase at the end of the metal strip. Since the transmittance is almost zero in all the minima (thus corresponding to complete destructive interference), we can conclude that the two modes are excited with a similar efficiency at the transition between the dielectric and the SPP waveguide and that they have a similar propagation loss along the SPP waveguide, as confirmed by the value of the imaginary part of the refractive index shown in red dashed lines in Fig. 4(a) and(b). From these figures it can also be appreciated that the difference between the real part of the effective index of the two SPP modes is almost independent of the thickness of the gold strip T and the width W (at least for W > 3 µm); this means that the beat length Δ𝐿𝐿 is not critically affected by fabrication tolerances. Moreover, the beat length does not change significantly with wavelength: Fig. 6 shows that the refractive index difference Δℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒 } = ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,0 } -ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,2 } between the two SPP modes changes by only ± 4% across a wavelength range of more than 60 nm, thus enabling the SPD to work on a broad wavelength range. This beating effect disappears for TE propagation, where SPP modes cannot be excited; in this case, Pout simply shows an exponential decay due to absorption in the metal strip. The shallow fluctuations on top of both TM and TE curves, with a period of about 500 nm, depend on Fabry-Perot (FP) cavity effects due to internal back reflections at the metal-dielectric interface; as the detector length L increases, this FP effect reduces because of the increased cavity loss due to metal absorption. As a result, the TE transmittance flattens after about 3 μm, while the TM mode maintains the beating effects between the two SPP modes. Numerical simulations were carried out to localize optical absorption across the waveguide. Since TiO2 is assumed to be fully transparent in the considered wavelength range (1550 nm), power dissipation occurs only within the Au layer. More in detail, in the cross-sectional plane the dissipated power is mostly concentrated within two-thirds of the core width [Fig. 7(a)], while it longitudinally spreads all along the detector, with a maximum dissipation at its center [Fig. 7(b)].

The absorption efficiency

𝜂𝜂 = 𝑃𝑃 𝑑𝑑 𝑃𝑃 𝑙𝑙𝑙𝑙𝑙𝑙𝑙𝑙 , ( 3
)
which is defined as the ratio between the power dissipated in the gold Pd and the overall losses Ploss, gives a direct measurement of the capacity of the gold to transform the lost power into heat. Note that Ploss includes the loss due to radiated light and back reflection due to mode mismatch at the input and output sections of the SPP waveguide. In Fig. 7(c), the simulated losses efficiency is shown, highlighting how the absorption of the TM mode (orange) exhibits both the oscillations mentioned previously: the fast fluctuations due to the Fabry-Perot cavity are superimposed on the wide oscillations due to the coupling of the SPP propagating modes. For short lengths, Ploss is dominated by the coupling mismatch and η is only about 10%. As the length increases the coupling losses remain constant but Pd increases and so does η. An efficiency of 42.5% is achieved for the TM mode at L = 2.3 µm. Notably in this working point, η is 1.7 times bigger than the value observed for TE propagation (about 25%), demonstrating the benefit of SPP propagation. The absorption efficiency increases up to 50.4% at L = 4.7 µm, which coincides with the second transmittance peak. This is a remarkable result since there are lengths for which the transmitted optical power and gold absorptance are maximum. The absorption efficiency is limited by the coupling losses and the mismatch between the waveguide mode and the SPPs modes, and it can be further improved by introducing a mode adapter at the input/output sections of the plasmonic waveguide.

B. THERMO-ELECTRICAL ANALYSIS

In the proposed photothermal detector, the optical power dissipated by the metal of the SPP waveguide causes a temperature variation that induces the bolometric effect [START_REF] Wu | Ultracompact on-chip photothermal power monitor based on silicon hybrid plasmonic waveguides[END_REF]. By measuring the change of the voltage drop due to the variation of the metal's resistance, it is possible to measure the optical power in the waveguide.

Figure 8 shows a top-view schematic of the complete device. The metal structure of the SPD is designed in order to enable the measurement of the light-induced temperature variation of the metal strip by means of a 4-probe technique. A small current Is is forced between the two pads on the left, while a lightdependent voltage change ΔV is measured with a highimpedance amplifier across the two pads on the right. The 4probes measurement allows to minimize the effect of the parasitic series resistance of the bonding wires and pads. Furthermore, a lock-in readout scheme is preferred for the readout, allowing to minimize the contribution of the 1/f noise of the preamplifier, maximizing in this way the sensitivity of the power monitor. For the design optimization of the electronic access to the detector, we varied the width of the sensing track Wtr from 200 nm to 1600 nm. A narrow width Wtr increases the equivalent thermal resistance from the plasmonic waveguide to the contact pads in such a way that, even though gold is a good thermal conductor, heat does not spread through the sensing tracks and remains mostly localized around the waveguide. Therefore, the smaller Wtr, the larger is the thermal resistance Rth, and hence the temperature variation. The temperature profile along the detector is presented in Fig. 9 for an input TM mode with a power Pin = 1 mW, reference room temperature (that is 25°C for Pin = 0). We considered two different lengths L of the detector and two different widths of the arms Wtr. For Wtr = 200 nm and L = 1.6 µm, we obtain a peak around 40°C, which is 15 °C higher than room temperature. These results also provide an optimal length for the sensing tracks as no temperature variation is observed after 10 µm (Wtr = 200 nm) or 20 µm (Wtr = 1600 nm). The thermoelectric performance of the proposed plasmonic detector is characterized by its responsivity Rd, defined as

𝑅𝑅 𝑑𝑑 = ΔV 𝑃𝑃 𝑑𝑑 , ( 4
)
where ΔV is the voltage difference observed when a current Is is forced through the circuit, and

Δ𝑉𝑉 = 𝐼𝐼 𝑠𝑠 Δ𝑅𝑅 = 𝐼𝐼 𝑠𝑠 𝑅𝑅 0 𝛼𝛼 Δ𝑇𝑇, (5)
where R0 is the resistance of the detector at reference room temperature and α is the temperature coefficient of the resistance of the gold.

Figure 10 shows the simulated voltage variation ΔV versus the optical power Pin at the input section of the detector. A temperature coefficient resistance α = 3.4⋅10 -3 [1/K] is considered. Simulations are performed when Is = 25 µA for Wtr = 200 nm (solid lines) and 1600 nm (dashed lines), and L = 1.6 µm (blue) and 2.3 µm (orange). Since the temperature varies linearly with Pin, and the resistivity varies linearly with temperature, ΔV is directly proportional to Pin. The voltage variation is very sensitive to the track width, for which we have a factor of 5 between the slope of the curves with Wtr = 200 nm and 1600 nm; much less sensitivity is observed versus the detector length L, where the slope change is less than 15%. This behavior is due to the fact that the total electrical resistance is dominated by Wtr rather than by L. Since the absorbed optical power Pd increases with L, the best detector is not that with the smallest footprint, but that with the best compromise between the absorption efficiency η and the total electrical resistance, which is L = 2.3 µm and Wtr = 200 nm. Fig. 10. Simulated variation of the voltage drop across the SPD for increasing power in the waveguide Pin for Is = 25 µA. The detectors with Wtr = 200 nm perform around 5 times better than the ones with Wtr = 1600 nm. Fixing the width of the tracks, the longer detector performs better, given the longer absorbing region.

III. TECHNOLOGY AND DEVICE FABRICATION

All devices are fabricated on a 22 mm by 22 mm silicon chip covered by a 2 µm thick thermally grown SiO2 layer. The devices are produced by three consecutive steps of electronbeam lithography. In the first step, the plasmonic gold detectors and their 4-probe addressing electrodes are patterned in a double-layer of PMMA (poly(methyl methacrylate)). This initial step is finalized by subsequent evaporation of a 3 nm Cr layer and a sub-40 nm Au film followed by a lift-off process. A 370 nm amorphous TiO2 layer is then deposited on the entire chip by electron-beam evaporation of a suitable target material on a sample holder maintained at 470 °C. Ellipsometry of this layer gives a refractive index of 2.2 at 1550 nm. The TiO2 PIC [START_REF] Andrieux | Optimized ICPCVD-based TiO2 photonics[END_REF] including waveguides and grating couplers is patterned in a second electron-beam lithography step in a negative tone resist which is then used as a hard mask during the following inductively coupled plasma reactive ion etching (ICP-RIE) of the TiO2 layer. Here, we use a CF4 plasma (20sccm) with a radiofrequency power of 50W, and an ICP power of 200 W, under a pressure of 4.5 mTorr at 0 °C to etch the layer. The last electron-beam lithography step consists of realizing macroscopic pads for electrically contacting the device with an external control and measurement electronics. The pads are made of 10 nm thick Cr and 200 nm thick Au layers to insure mechanical robustness against wire bonding. The TiO2 waveguides are 1200 nm-wide, 370 nm-high and 1.5 mm-long. A portion of such a waveguide and a dummy Au detector (without the sensing tracks) are shown in Fig. 11(a). In this

IV. EXPERIMENTAL RESULTS

Light coupling into the chip was achieved by using vertically emitting grating couplers designed to operate on TM polarization. Straight waveguides integrating SPDs with increasing length L were fabricated to assess the insertion loss (IL) and the back reflections of the SPD. Fig. 12 shows the transmitted optical power Ptx versus L normalized to a reference all-dielectric waveguide (L= 0). Experimental data are averaged over a wavelength range from 1520 nm to 1580 nm, where the SPDs exhibit a flat spectral response, and the error bars indicate the standard deviation evaluated on several nominal identical samples. A very good is found with numerical simulations (data from Fig. 5 are reported in dB-scale for a direct comparison) on the position of the minimum transmission (around 1.1 µm) and the maximum transmission (2.2 µm). The insertion loss in the transmission maxima (5.5 dB) is higher than the value expected from simulations (2.5 dB); we believe that this is caused by some scattering loss in the plasmonic waveguide due to surface roughness at the metaldielectric interface. The back reflection of each SPD was measured by using coherent optical frequency domain reflectometry (OFDR) [START_REF] Morichetti | Roughness Induced Backscattering in Optical Silicon Waveguides[END_REF] and amounts to -17 (± 0.7) dB, in agreement with electromagnetic simulations. This means that mode adapters from the dielectric to the plasmonic waveguide are not necessary, even though they could reduce the scattering loss.

Fig. 13 shows the light-dependent change of the voltage ΔV across the SPD, as measured by a lock-in scheme, for increasing power Pin in the waveguide, when a current Is = 25 μA is applied. Is is constant and does not change when light travels through the waveguide, nor when the optical power gets dissipated in the gold. Also, being 25 µA and the optical power inside the WG > -20 dBm, the dissipated power due to the light is 2 orders of magnitude higher than that due to Is. For these reasons, any possible thermal effect caused by Is is negligible. In agreement with the numerical results of Fig. 10, ΔV increases linearly versus Pin. The responsivity Rd, which is given by the slope of the linear fit (red), is equal to 7.5 μV/mW in good agreement with simulations of Fig. 10. With the lock-in amplifier bandwidth set to 1 Hz, the minimum detectable power is about -20 dBm, which corresponds to a voltage change ΔV = 75 nV, that is induced by a temperature increase by only 1.75e-3 °C. The detector sensitivity can be extended to lower optical power by increasing Is, the limit being given by the current handling capability of the thin arms of the detector, which in our case is about 250 µA. With this supply current we were able to measure a minimum optical power of -30 dBm in the waveguide. In the considered device the dynamic range is restricted to 15 dB by the maximum power in the waveguide (Pin = 0.75 mW), which is limited by the coupling efficiency of the grating couplers (13 dB loss each). The time response of the SPD is shown in Fig. 14. It was measured by modulating the input light with a square wave at 20 kHz and reading the detector voltage ΔV with an oscilloscope triggered by the modulating signal. The time constant of the detector response (red curve), which is defined as the time it takes for the system step response to reach 1 -e -1 ≈ 63.2% of its final value, is about 1.2 µs which corresponds to a bandwidth of about 132 kHz. The measured response is in very good agreement with simulations (yellow curve) performed by using a time-dependent FEM model, showing a time response of 1.1 µs. The limit to the SPD speed is set by the heat capacity of the dielectrics and the gold's resistivity thermal constant (the electronics response in blue). Fig. 14. Normalized time response of the SPD. In blue, the response time of the SPD when electronically and in orange when optically excited. In yellow, the simulated response time.

V. CONCLUSION

In conclusion, we demonstrated the possibility of monitoring the light intensity in TiO2 waveguides by exploiting photothermal effect in a multimode SPD. By optimising the geometry of the device, we exploited constructive interference of the SPP modes in order to minimize the coupling loss at the dielectric-plasmonic waveguide interface and optimize the absorption efficiency in the metal. This allows an absorption efficiency, and hence a photothermal response, that is 1.7 times bigger than the value observed for TE propagation, where SPP propagation is inhibited. As a result, the fabricated device, which is extremely compact in size (1.6-μm-long), exhibits a sensitivity (-20 dBm) and a time response (>100 kHz) that fulfils the requirements for implementing monitor and control operations in PICs. The SPD insertion loss, which is about 5 dB in the fabricated devices (2 dB in simulations) can be substantially reduced by introducing a mode adapter at the input/output sections of the plasmonic waveguide [START_REF] Shi | Enhanced coupling efficiency between dielectric and hybrid plasmonic waveguides[END_REF]. The fabrication process is extremely simple and versatile since it can be ported to any dielectric waveguide platform, where monolithic integration of semiconductor photodetectors is not a viable approach.

VI. ACKNOWLEDGMENT

This work was partially performed at Polifab, the micro-and nanofabrication facility of Politecnico di Milano (www.polifab.polimi.it).

Fig. 1 .

 1 Fig. 1. Front (a) and lateral (b) view of the platform with the SPD. The width and thickness of the TiO2 core are 1.2 µm and 370 nm, respectively. A thin strip of gold is realized below the TiO2 core with thickness T, width W, and length L, which are the SPD parameters optimized by numerical simulations.

Fig. 2 .

 2 Fig. 2. Propagation of a surface plasmon excited by a TM dielectric mode. When the TM mode encounters the metal strip (yellow), it may excite several SPP modes (indicated with different colors) that propagate at different phase velocities and interfere constructively or destructively at the end of the plasmonic waveguide depending on the strip length L.

Fig. 3 .

 3 Fig. 3. Electric field profile of the guided modes of the plasmonic waveguide with T = 34 nm and W = 3 mm: (a) first order TM0 and (b) third order TM2 SPP modes. The second-order mode (not shown) is not excited by the TM mode of the input dielectric waveguide and does not contribute to the photo-thermal response of the device.

Figure 4 (

 4 Figure 4(a) shows the real part (solid blue line) and the imaginary part (solid orange line) of the effective index of the fundamental mode versus the thickness T of the metal strip. The propagation loss

Fig. 4 .

 4 Fig. 4. (a) Real (blue) and imaginary (orange) parts of the SPP modes refractive index versus the thickness T of the gold strip. (b) Real (blue) and imaginary (orange) parts of the SPP modes refractive index vs gold width.

  (b) we show the electric field profile of the third mode, whose complex effective index is shown in Figs 4(a) and (b) in dashed lines. The second mode is not shown because its electric field has an odd symmetry in the horizontal direction and it is not excited by the fundamental TM mode of the dielectric waveguide. The real part of the refractive index of the first and third modes is substantially different, this meaning that the two modes travel with a different phase velocity and accumulate a relative phase shift along the SPP waveguide. When the relative phase difference between the two SPP modes is a multiple of 2π, the overlap with the fundamental TM mode of the output dielectric waveguide is maximized and Pout is maximum. The beat length Δ𝐿𝐿 is related to the difference between the real part effective index of the two SPP modes as Δ𝐿𝐿 = 𝜆𝜆 0 ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,0 }-ℜ{𝑛𝑛 𝑒𝑒𝑒𝑒𝑒𝑒,2 }

Fig. 5 .

 5 Fig. 5. Normalized transmittance through the SPD versus the metal strip length L for TE (blue) and TM (orange) input modes. The width and thickness of the strip are W = 3 µm and T = 34 nm.

Fig. 6 .

 6 Fig. 6. Simulation of the difference between the real part of the effective refractive index of the first SPP mode (blue) versus wavelength. The simulation suggests that the difference between the real part of the modes remains constant over a large bandwidth.

Fig. 7 .

 7 Fig. 7. Front (a) and lateral (b) view of the simulated power dissipated density. (c) Absorbance efficiency η, defined as the ratio between the power dissipated in the gold strip and the overall loss of the SPD versus the detector length L for TE (blue) and TM (orange) modes.

Fig. 8 .

 8 Fig. 8. Schematic of the SPD.

Fig. 9 .

 9 Fig. 9. Transversal SPD temperature profile for an input power Pin = 1 mW.

  length L of the detector placed under the waveguide is 1.6 µm by a width W of 3 µm. Figure11(b)shows an overview of a detector fully connected to its four-probe electrode systems.

Fig. 11 .

 11 Fig. 11. False color SEM photographs of the TiO2 testing PICs. (a) Dummy detector L = 1.6 µm and W = 3 µm placed underneath a TiO2 waveguide. (b) overview of the SPD and its addressing electrodes arranged in a 4-probe configuration. The yellow hue highlights the Au parts of the device.

Fig. 12 .

 12 Fig. 12. Normalized transmitted power vs. SPD length for TM mode. The experimental curve (blue) is in good agreement with FEM simulation (orange).

Fig. 13 .

 13 Fig. 13. Measured responsivity of the SPD for Is = 25 µA. The voltage is linear w.r.t the optical power and the responsivity equals 7.5 μV/mW

This work was partially supported by the European Commission through the H2020 project Nebula (Contract No. 871658).


Nhat Jaysen Nelayah

Tai Nguyen

Damien Alloyeau

Guillaume Yangshu

Christian Ricol- Leau

Long-range chemical orders in Au-Pd nanoparticles revealed by aberration-corrected electron microscopy †

Despite the importance of gold-palladium nanoalloys in heterogeneous catalysis, the phase stability of Au-Pd alloy still remains unclear. We report here on the alloying and chemical ordering in epitaxially-grown and post-annealed gold-palladium nanoparticles (NPs) using aberration-corrected transmission electron microscopy. Au-Pd NPs with controlled size, composition and structure, were grown by pulsed laser deposition on freshly-cleaved NaCl(001) single crystals heated at 300 • C. After transfer to an amorphous carbon support, the NPs were annealed in vacuum at elevated temperatures above 400 • C for few hours (6-10 hours) to promote chemical ordering. The as-grown NPs were mostly monocrystalline with a chemically-disordered face-centered cubic structure. Upon high-temperature annealing, high degree of chemical ordering was observed in nanometer-sized NPs. Electron microscopy measurements showed that both L1 0 and L1 2 orders are stabilized in the Au-rich region of Au-Pd phase diagram. These ordered phases exist at temperatures as high as 600 • C. Moreover, compositional analysis of single annealed particles revealed that the observed chemical ordering occurs in parallel to a two-tiered Ostwald ripening process. Due to this ripening process, a clear dependence between chemical composition and particle size is established during annealing with an enrichment in Pd as the NPs grow in size. Our results, besides clarifying some controversial aspects about long-range order in Au-Pd alloy, sheds light on the structural stability of Au-Pd nanoalloys at elevated temperatures.

Introduction

Gold-based bimetallic nanoparticles (NPs) are presently at the center of much scientific and technological interests due to their unique thermodynamic, optical and catalytic properties. In the field of heterogenous catalysis, among the wide range of gold-based bimetallic alloys, gold-palladium NPs have been widely studied. This is due primarily to Au-Pd alloys being an active catalysts in various reactions (in the synthesis of vinyl acetate monomer 1 , selective oxidation of primary carbonhydrogen bonds in Toluene 2 , oxidation of primary alcohols to aldehydes 3 , direct synthesis of hydrogen peroxide by reacting molecular hydrogen and molecular oxygen at ambient temperature 4 ). In these reactions, Au-Pd NPs exhibit enhanced activity, selectivity and stability when compared to either pure Pd or Au nanocatalysts. The promotion of the catalytic properties in nanoalloys is explained on the basis of an enhancement of both electronic (ligands effects) and geometric properties (ensemble effects) upon intermetallic mixing 5 . In Au-Pd nanostructures, the ensemble effect is often assigned to a dilution † Electronic Supplementary Information (ESI) available: [details of any supplementary information available should be included here]. See DOI: 10.1039/b000000x/ Laboratoire de Matériaux et Phénomènes Quantiques, Bâtiment Condorcet, CNRS UMR 7162,Université Paris Diderot, Paris, France. Fax: 01 5727 6241; Tel: 01 5727 6998; E-mail: jaysen.nelayah@univ-paris-diderot.fr of Pd surface atoms by gold. For instance, in the oxidation of glycerol to glyceric acid, a high activity and selectivity is associated with isolated Pd monomer sites in contact with Au atoms 6 . For the rational design of Au-Pd nanocatalysts, chemical and physical approaches have been developed. Chemical approaches involving colloidal synthesis methods (such as co-reduction and seeded-growth approaches) have yielded both completely random alloys 7,8 and core-shell structures 9 . Production of bimetallic Au-Pd nanoparticles/clusters has also been achieved by vapor deposition methods 10,11 .

Information about chemical order and phase stability in Au-Pd alloy have been compiled by Okamoto and Massalski 12 . According to the experimental bulk phase diagram, Au and Pd show complete miscibility at all compositions in the solid phase. The resulting solid solution displays a random mixing arrangement with a disordered face-centered cubic (fcc) structure. In the fcc lattice, each atomic site is occupied by either a gold or palladium atom. The occupation probability of a given site depends on the ratio of the constituent elements. The most surprising features of this phase diagram is the presence of L1 0 -and L1 2 -types long-range ordered phases around compositions AuPd(L1 0 ), Au 3 Pd(L1 2 ), and AuPd 3 (L1 2 ) despite that no chemical order has ever been observed experimentally in bulk Au-Pd. Inclusion of ordered phases in the bulk phase diagram is based on the assessment, by electron diffraction anal-

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ysis, of the atomic structure of Au-Pd thin films evaporated on cleaved surfaces of KCl and NaCl salts and post-annealed at high temperatures [13][14][15][16] . Around compositions 25 and 75 atomic per cent (at.%) Au, a L1 2 -type ordered structures have been experimentally observed with order-disorder (O/D) transition temperatures estimated at 850 • C at Au 0.8 Pd 0.2 14,16 and between 780 and 900 • C at Au 0.2 Pd 0.8 composition 15 . Around the equiatomic composition, a L1 0 -type order was identified by Nagasawa in thin AuPd film annealed at 500 • C during 4 hours 13 . The corresponding O/D transition was observed between 650 and 700 • C. The L1 0 structure is based on a tetragonal lattice characterized by planes of pure Au atoms and pure Pd atoms alternately stacked along the c-axis of the tetragonal unit cell. As for the L1 2 order, it occurs on a fcc lattice with either Au (in AuPd 3 ) or Pd (in Au 3 Pd) occupying the corners of the cube and the other element centered on the six faces of the cubic unit cell. The unit cells (Fig. 1S) and space groups of the fcc, L1 0 and L1 2 structures are included in the Electronic Supplementary Information (ESI). It should be noted that the existence of this L1 0 phase is still uncertain. Firstly, in a later paper 14 , Nagasawa questioned the reliability of his experiments showing the existence of L1 0 phase. Secondly, based on short-range order parameter measurements in Au-Pd alloys, the temperature range of stability of the L1 0 order is expected to be lower than 100 • C 17 . Recently, a L1 1 order has been observed experimentally in core-shell NPs with a Au octahedron core and a Pd epitaxial shell heat-treated at 250 • C 18 . This L1 1 order was limited to the interfacial region between the core and the shell and was assigned to the diffusion of Au atoms from the core to the shell so as to release interface mismatch. In this study, no other chemical order phase was identified.

Gaining deeper insights into the phase stability of Au-Pd alloys has also relied intensively on atomistic simulations. The equilibrium configurations of Au-Pd alloys have been studied by Monte Carlo simulations using embedded-atom method interatomic potentials 19,20 and Density Functional Theory calculations combined with cluster expansion method 21,22 . At equiatomic composition, the low-temperature ordered phases derived from first-principles cluster expansion methods, by Barabash et al. 21 and Sluiter et al. 22 , predict bulk AuPd to be stable in chalcopyritelite Au 2 Pd 2 superlattice structure 23,24 . In comparison, Monte Carlo simulations predict the existence of L1 0 and L1 2 orders in both Au-Pd bulk alloy and nanoalloys (consisting of about 1000 atoms or less) with cuboctahedral, truncated octahedral and spherical shapes 19 . In these nanoalloys, chemical order is predicted in the core alongside systematic surface and sub-surface enrichments in Au and Pd respectively. In bulk AuPd alloy, the L1 0 order is calculated below 100 • C. As observed experimentally in CoPt nanoclusters 25 , a depression of the O/D temperature is expected in nanoclusters.

Developing new Au-Pd catalysts or enhancing the perfor-mances of existing ones on a rational basis would gain significantly from a clear knowledge of the phase diagram of Au-Pd NPs. Due to the scarcity and reliability of existing experimental and calculation data, the existence of chemical order in Au-Pd alloy is still a highly debating issue. In this paper, through the fabrication of Au-Pd nanoalloys and study of their structure by aberration-corrected transmission electron microscopy (TEM), long-range orders have been unambiguously identified in Au-Pd alloy. These results constitute a major step towards a complete understanding of the phase stability of Au-Pd alloy.

Experimental conditions

Gold-palladium NPs were synthesized by pulsed laser deposition (PLD) technique in a high vacuum chamber under a pressure of about 10 -7 Torr 26 . The Au-Pd NPs were deposited on a 1×1 cm 2 freshly cleaved sodium chloride NaCl(001) singlecrystal surface. To ensure well-oriented growth of the NPs on the substrate, the latter was heated at 300 • C during deposition.

Real-time monitoring of the temperature was achieved using Eurotherm controller and K-type thermocouple. The transfer of material from the target to the substrate was done by using a KrF excimer laser with wavelength of 248 nm with pulse duration of 25 ns and a repetition rate of 5 Hz (COMP-exPro 102 F, COHERENT R ). Owing to the ablation thresholds of Au and Pd, the pulse energy at the exit of the laser was set in the range 250-400 mJ. To fabricate Au-Pd nanoalloys, two ultra pure (99.99% purity) Au and Pd targets were alternately ablated. To ensure complete mixing of the constituent atoms, the deposition was divided in sequences. In each sequence, the total thickness of deposited metals was fixed to 1 Å. The deposition rate for each metal was controlled by an in situ quartz balance prior to deposition. As the constituent metals can be mixed in any ratio, our set-up enabled a precise control of both the composition and nominal thickness of alloy deposited (which corresponds to the total thickness of metals deposited in a continuous film approximation). In the present study, the nominal thickness deposited was fixed to 1 nm to ensure the formation of well-separated NPs with sizes of the order of a few nm. After deposition, the Au-Pd NPs were covered by a few nm-thick amorphous (a-)carbon film obtained by ablating a pure carbon target present in the same set-up. Carbon replica was obtained by dissolving the sodium chloride in desionized water followed by the transfer of the carbon-supported NPs to standard copper TEM grids (holey carbon on a 400 mesh Cu grid, Agar Scientific) for post-annealing experiments and structural investigations by transmission electron microscopy (TEM). More information about the preparation of carbon replica can be found in 27 .

Prior to annealing and TEM analysis, the replicas were covered by another thin film of a-carbon to limit oxidation of the bimetallic NPs when the grids are exposed to air and sintering of NPs during thermal annealing at high temperatures. Postannealing of the a-carbon supported NPs was done at temperatures between 400 to 600 • C for different time intervals. Annealing experiments were undertaken in the growth chamber using the same heating stage as the one used to heat the rock salt substrate during deposition. After annealing, the specimen was quenched to room temperature in vacuum. The pressure in the chamber during annealing was kept around 10 -7

Torr. The structure of the as-grown and annealed NPs was studied by standard and aberration-corrected TEM at 200 kV. Low magnification imaging, selected-area electron diffraction (SAED) and energy dispersive X-ray spectroscopy on NP assemblies (TEM-EDS) were undertaken in a JEOL 2010 TEM.

All SAED patterns are obtained by selecting nanoparticles distributed over circular area with a diameter of about 560 nm. For ultra-high resolution TEM (UHRTEM) imaging and single-particle X-ray analysis in scanning mode (STEM-EDS), a JEOL ARM 200F microscope was used. This microscope combines a newly-developed cold field emission gun 28 and a CEOS hexapole spherical aberration corrector (CEOS GMBH) 29 to compensate for the spherical aberration of the objective lens. This unique TEM configuration achieves a point resolution of 75 pm at 200 kV 30 . High resolution images have been acquired with negative Cs values up to -500 nm and with the objective lens slightly overfocused. In these conditions, the image contrast due to the carbon matrix is minimized allowing particle imaging with high signal-to-noise ratio and atomic columns appear in white at the Scherzer defocus. STEM-EDS spectra were obtained with acquisition times of 50 or 100 s. A typical raw STEM-EDS spectrum is shown in Fig. 2S. Particle composition was determined by analyzing the intensities under the Au-M α,β (between 2.00 and 2.24 keV) and Pd-L α,β (between 2.70 and 2.94 keV) edges using the Cliff-Lorimer method 31 . The Cliff-Lorimer K factor used for quantification was experimentally determined by utilizing a thin wedge-shapes Au 48 Pd 52 film fabricated by PLD and thinned by focused ion beam milling. The composition of the film was precisely determined by inductively coupled plasma optical emission spectrometry prior to thinning. Instead of using a fixed probe, the electron beam was continuously scanned in a controlled manner over a rectangular region on the NPs during STEM-EDS measurements so as to minimize electron beam induced damage and ensure that the measured composition is as close as possible to the real one. For each particle, the scanned region was chosen to encompass as much as possible its projected area.

Results

Structure of faceted Au-Pd NPs observed by selected area electron diffraction and ultra high resolution transmission electron microscopy imaging

Au-Pd nanoparticles with composition between AuPd 3 and Au 3 Pd were fabricated. A low magnification bright field TEM image of an assembly of as-grown NPs on the a-carbon support is shown in Fig. 1. The NPs display mostly square or rectangular outlines. The average projected area diameter, i.e. the diameter of a perfect circle enclosing an area equivalent to the projection area of the NP, follows a narrow distribution with a mean diameter of 4.8 nm (see inset of Fig. 1a). According to TEM-EDS measurements (data not shown here), the mean composition of the present specimen is 60±5 at.% Au (using the X-ray detector fitted to the JEOL 2010 TEM). The SAED pattern, obtained from the assembly of NPs in the TEM image, is shown in Fig. 1b. It should be noted that all diffraction patterns presented in this work were acquired on NPs with sizes under 20 nm. The SAED pattern of the as-grown NPs is composed of well-defined diffraction spots. This single crystal-like electron diffraction pattern indicates the presence of well-oriented Au-Pd NPs on the carbon replica. The production of well-oriented NPs is a result of the epitaxial growth of the bimetallic nanoparticles on the rock salt substrate. The predominant epitaxial relationship between the metals and the salt substrate is most likely the cube-on-cube orientation epitaxy (100)Au-Pd (100)NaCl, [001]Au-Pd [001]NaCl as previously observed in epitaxially-grown Pd NPs on NaCl(001) [32][33][34][35] . The relative orientation of the fabricated NPs is preserved during the preparation of carbon replica. Indexing the SAED pattern shows that it is composed of sharp and intense 200 and 220 fundamental reflections. These reflections are assigned to the presence of completely alloyed Au-Pd NPs having disordered fcc structure and oriented such as one of the axes of their cubic lattice are perpendicular to the support. The as-grown NPs were also imaged by ultra-high resolution TEM (UHRTEM). Fig. 1c shows UHRTEM images of two as-grown NP in <001> orientation alongside the corresponding Fourier transforms (FT). The FT patterns show the 200 and 220 spots of the fcc structure. Both NPs are faceted single crystal with main edges and minor truncations parallel to the <110> and <100> directions respectively. As previously observed in Pd NPs evaporated on NaCl at 150 • C 32 , the shape of the Au-Pd NPs, dictated by the growth kinetics, is most probably a truncated octahedron. In the <001> orientation, the latter appears as a square/rectangle. In addition to single crystals, a few polycristalline fcc NPs are also observed. Their average size exceeds 10 nm. NPs with structure (composition, atomic arrangement) comparable to those presented in Fig. 1 were subjected to long high-temperature annealing in vacuum. Structural changes that occur during annealing are described in the next section. TEM images and the corresponding SAED patterns of two annealed specimen (A and B) are shown in Fig. 2. Sample A (Fig. 2a) consists of NPs, similar to those in Fig. 1, annealed at about 600 • C for 10 hours. Sample B (Fig. 2b) consists of NPs, with same composition as in Fig. 1, annealed at about 500 • C for 8 hours. It should be noted that the as-grown NPs annealed in sample B had an average in-plane size double that of the NPs shown in Fig. 1a (comparison of the histograms in Fig. 1 and3S). This explains the slightly larger in-plane size and lower surface coverage observed in sample B as com-pared to sample A. From the comparison of Fig. 1 and2, two observations can be made about the influence of annealing. Firstly, a clear evolution of the nanoparticle's in-plane morphology of the NPs from square/rectangular to rounded profiles is observed upon annealing (Fig. 2a-b). Secondly, the mean in-plane size is increased during annealing in both samples. In sample A, analysis of 240 NPs from the TEM image in Fig. 2a gives a projected area diameter equals to 5.9 ± 1.9 nm (as compared to a mean in-plane size of 4.8 nm before annealing). Similarly, the in-plane diameter measured in sample B is 8.5 ± 2.1 nm as compared to 6.5 ± 2.2 nm before annealing (Fig. 3S). The SAED pattern of sample A is of powder-type (Fig. 2c). This results from the loss, during annealing, of the preferential <001> orientation perpendicular to the support in the as-grown sample. On annealing at 600 • C, we note in the SAED pattern (i) a splitting the 002 diffraction rings indicating a loss of the original cubic structure, (ii) the appearances of 001 and 110 superlattice reflections and (iii) the absence of the 100/010 superlattice reflection. The size distribution in annealed sample B is comparable to that observed in A (see insets of Fig. 2a-b for a comparison of the size distributions of the annealed NPs in A and B). The SAED pattern of the NPs in the TEM image shown in Fig. 2d shows 001 and 110 superlattice reflections indicating that chemical order is again present. But, no splitting of the 002 reflections appears here. The SAED results from both samples are consistent with the presence of chemical orders of either tetragonal L1 0 -type or cubic L1 2 -type. The splitting of the 002 diffraction rings observed in sample is a strong indication of the predominance of the L1 0 order in that case. To determine unambiguously the chemical orders that exist in samples A and B, the structure of the annealed NPs was studied on single particle level by UHRTEM imaging. Fig. 3 shows UHRTEM images of chemically-ordered NPs from annealed samples A and B together with the corresponding FT images (below each UHRTEM image). These measurements show that the chemical order in sample A is predominantly of L1 0 -type. Fig. 3a show a L1 0 NP imaged close to the [100] orientation. In this orientation, the tetragonal caxis is in the plane of the carbon support. Analysis of the corresponding FT spectrum shows the presence of 001 superlattice reflections. In Fig. 3b-c, the L1 0 -ordered NP is close to the [001] orientation, i.e. with its quadratic c-axis aligned normal to the carbon substrate. 110 superlattice reflections are clearly visible. The appearance of both the 001 and 110 superlattice reflections in the SAED pattern shown in Fig. 2c results from the presence of all three variants of the L1 0 phase in the annealed sample. These variants differ by the orientation of the c-axis of the tetragonal lattice with respect to the plane of the substrate. The value of the tetragonality ratio (c/a ratio) of the L1 0 order, deduced from the previously shown SAED data and UHRTEM images of NPs close to the [001] orientation, is about 0.92. This value is very close to the one measured experimentally by Nagasawa in AuPd thin films (c/a = 0.93) 13 . In sample B, the ordered NPs are mostly of L1 2type. Fig. 3d-f shows UHRTEM images of L1 2 -ordered NPs from sample B. In Fig. 3d, the particle is image close to the [001] orientation. L1 2 -order is evidenced by the presence of 100 and 110 superlattice reflections in the FT pattern. The NPs in Fig. 3e-f are imaged along the [101] and [112] orientations respectively. Again, 110 superlattice reflections are seen in the corresponding FT patterns. High-resolution measurements such as the one presented here show that the ordered NPs are mostly single crystals with truncations along the <110> and <100> directions. In comparison to as-grown NPs, truncations along the two sets of direction are equally favorable in the annealed NPs and result in their rounded morphologies. Finally, it should be noted that the results present here are extracted from a systematic study of the thermal stability of Au-Pd nanoalloys as a function of annealing temperatures and durations. In this study, particle composition was between the AuPd 3 and Au 3 Pd. Annealing temperature was set between 500 and 700 • C for duration up to 24 hours. Our annealing experiments showed that no chemical order is stabilized when the annealing time was less than 6 hours irrespective of particle composition and annealing temperature

Atomic ordering observed by selected area electron diffraction and ultra high resolution transmission electron microscopy imaging

Discussions

In this study, L1 0 and L1 2 chemical orders have been stabilized in Au-Pd nanoalloys. To access the stable phase of Au-Pd alloy in the composition-temperature space, the NPs with initial completely disordered fcc structure are annealed at high-temperatures. In this work, loss of structural uniformity is a common consequence of annealing. Intriguingly, subtle changes in annealing parameters lead to drastic changes in the ordering processes that occur in Au-Pd nanoalloys during high-temperature annealing. Indeed, annealing at a temperature of 600 • C for 10 hours results in a L1 0 order (sample A) while decreasing the annealing temperature to about 500 • C leads to a L1 2 order after 8 hours (sample B). As it will be described in the following, this experimentally-observed tunability of the chemical order via annealing conditions can be explained on the basis of a change in the particle's composition by thermally activated mechanisms of Ostwald ripening during annealing processes.

The chemical composition of the annealed NPs from both samples were measured by single-particle X-ray analysis. As described in the ESI, the errors in composition measurements by STEM-EDS were between 1 and 3 %. To determine the size of every particle analyzed, particle size was determined precisely by acquiring a dark-field image in scanning mode prior to X-ray analysis (see Fig. 2S). The dark-field image Fig. 4 The elemental composition of the NPs in sample A ( ) and sample B (•). Measurements of the composition in single particles were done by STEM-EDS. In the annealed samples, the elemental composition is clearly size-dependent with an enrichment in Pd as the NPs grow in size. This composition-size dependence is due to the activation of Ostwald ripening processes. In comparison, little variation of the composition from particle to particle is observed in the as-grown sample (60-70 at.% Au). as shown by the STEM-EDS analysis of 12 as-grown NPs from the sample shown in Fig. 1a ( ). was then analyzed with the particle size analysis tool in Digital micrograph software to yield particle size with an error of 0.4 nm. Fig. 4 displays the variation of the elemental composition of annealed NPs in samples A ( ) and B (•) as a function of particle's in-plane diameter up to about 85 nm. For comparison, the elemental compositions of 12 as-grown NPs from the sample shown in Fig. 1 are also presented ( ). The inset of Fig. 4 shows a zoom-in view of the size-composition dependence for diameter below 20 nm. In the as-grown sample, STEM-EDS measurements show that the compositions of the NPs are between 60 and 70 at.% Au and are slightly size-dependent. This size-composition dependence can result from Ostwald ripening activated during particle growth. Upon annealing at 600 • C in vacuum and in parallel to particle sintering, the composition of the annealed NPs are found to vary significantly with particle size. A clear enrichment in Pd is observed as the NPs grow in size. When the diameter is lower than about 5 nm, the NPs are nearly monometallic Au structures while for diameters exceeding 50 nm, the particle's composition tends towards Au 30 Pd 70 . Between these extreme compositions, the Au content decreases monotonously with size. A similar trend is obtained when analyzing sample B. However, due to the lower annealing temperature, particle sintering is less severe in this sample. Indeed, the particle size does not exceed 60 nm where the Au content is about 40 %. The simultaneous evolution of particle size and com-

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position observed during high-temperature annealing can be understood by considering the processes that contribute to particle sintering.

Particle sintering in nanoalloy occurs through two main mechanisms : substrate mediated particle coalescence (static and dynamic) and Ostwald ripening process. During static coalescence, atomic diffusion takes place between immobile NPs via the support or the medium encapsulating the NPs. As for dynamic coalescence mechanism, it involves NPs that diffuse on the support and when in close interaction, merge to form bigger NPs. Coalescence mechanisms do not give rise to any variation in particle composition as sintering proceeds. As for Ostwald ripening, it is a thermoactivated process. During Ostwald ripening, large thermodynamically favorable NPs grow by taking up mobile atoms at the expense of smaller ones which shrink. In the case of an assembly of monometallic NPs, Ostwald ripening results solely in a broadening of the size distribution. When dealing with NPs made of two different metals, Ostwald ripening can lead to a sizedependent composition as shown in CoPt 36 , IrPd 37 and AuPd 38,39 nanoalloys. In CoPt subject to high-temperature annealing, an enrichment in Co as the particles grow in size was observed. This compositional change has been attributed to a two-tiered Ostwald ripening mechanisms with respect to the transfer of Co and Pt atoms between particles. The evaporation rate of Co atoms from nanoparticle being higher than that of Pt atoms leads to an enrichment in Co with particle size.

Ostwald ripening in carbon-supported Au-Pd nanoclusters has been observed both at room temperature upon exposure to hydrogen gas 39 and during calcination at 400 • C 38 . Hydrogeninduced Ostwald ripening occurs since in the presence of hydrogen, the binding energy of the constituent atoms of a Au-Pd NP is reduced leading to their spontaneous detachment from the NP. Since the detached Au and Pd atoms have different mobility on the TEM carbon support, the composition of the NP varies as Ostwald ripening proceeds, with the faster Pd atoms enriching the large NP. In the work of Herzing et al where Ostwald ripening was induced during calcination,the size-composition dependence was opposite to the one observed in the presence of hydrogen by Di Vece et al. Indeed, the latter showed that the Au-to-Pd ratio decreased with NP size. In the present work, Ostwald ripening is thermally induced in vacuum and the composition variations observed can be fully understood by considering the energetics of the two limiting mechanisms associated with this ripening process: detachment of the metal atoms from the clusters and diffusion of the detached atoms until they are captured by other clusters on the support. The detachment of atoms from a cluster is governed by the cohesive energy E c of the cluster and the energy of adsorption E a of the detached atom on the surface. The amount of energy required to detach an atom from a cluster is equal to the energy difference between its final (on the support) and initial states (in the cluster), i.e. E a -E c . Thus, to activate the evaporation of atoms from a cluster, the minimum energy required is equal to the energy barrier (E a -E c ). The experimental cohesive energy per atom for bulk gold and palladium are 3.81 eV and 3.89 eV respectively [START_REF] Kittel | Introduction to Solid State Physics[END_REF] . It should be noted that these large values of the cohesive energy exclude reevaporation of metal atoms in the absence of any thermal excitation. The absorption of gold and palladium atoms have been studied on graphene by first-principles calculations [START_REF] Liu | [END_REF] . Estimation of adsorption energies per atom for Au and Pd adatoms on graphene are 0.10 and 1.15 eV respectively. These values leads to an evaporation barrier per atom equals to 3.71 eV and 2.74 eV for Au and Pd evaporation respectively. These energies are much higher than the diffusion barriers in noble metals which are very small with only few meV. Thus, Ostwald ripening in Au-Pd alloys is limited by the evaporation barriers which favor the detachment of Pd atoms. This preferential transfer of palladium atoms between NPs accounts for the size-composition dependence measured experimentally. The growth kinetics and evolution of composition of the annealed NPs are influenced by both the annealing conditions and the initial structure (composition, size) of the as-grown NPs.

In the sample A, a large spread out of the composition is observed in the size domain 5-20 nm (see inset of Fig. 4). For instance, for a particle diameter around 10 nm, the composition varies between 60 at.% Au and that of pure gold. From these data, we cannot exclude the presence of L1 2 order in sample A. Indeed, our UHRTEM measurements shows that few NPs with L1 2 -order were present in sample A in the size range below 20 nm. In sample B, variations in composition are less significant in the size domain 5-20 nm. The Au content remains mainly between 80 and 90 at.% Au. At these compositions, the predominant order is expected to be of L1 2 -type (the order parameter being the highest at Au 3 Pd composition). This is confirmed by UHRTEM analysis (as those shown in Fig. 3) of a large number of NPs from sample B. It should be noted that few L1 0 NPs with sizes below 20 nm were also observed in our high-resolution measurements. The composition of these NPs should be close to the equiatomic composition at which the L1 0 order is highest. The formation of these NPs is again a result of the thermally-activated Oswald ripening during annealing. Finally, due to the large distribution in composition after annealing, the ordered NPs were found to coexist with large number of chemically-disordered fcc NPs in both samples A and B.

Conclusions

In this work, epitaxially-grown Au-Pd nanoalloys were annealed in vacuum at temperatures above 400 • C for few hours (6-10 hours). The initial structure of the nanoalloys was of disordered fcc-type. After annealing, L1 0 -and L1 2 -chemically-ordered monocristalline NPs were observed in the Au-rich domain of the phase diagram at temperatures between 500-600 • C. Single particle X-ray analysis showed that chemical ordering occurred in parallel to a two-tiered Ostwald ripening activated during high-temperature annealing. As a result of this thermally-activated Ostwald ripening, the particle size distribution is inevitably broadened and a size-composition dependence is established upon annealing. The composition distribution of the latter is found to be highly dependent on the annealing conditions with the general trend being an enrichment in Pd as the NPs grow in size. These results constitute the first experimental evidence of the stabilization of longrange orders in Au-Pd nanoalloys. In terms of their catalytic properties, these ordered Au-Pd NPs may offer a new class of advanced nanocatalysts for various chemical reactions. However, in order to fully realize the potentials of these ordered nanostructures in heterogeneous catalysis, the detrimental effects of particle sintering at high temperatures should be limited and fully understood. Moreover, we strongly believe that both ab initio and semi-empirical simulations of the thermodynamic properties of Au-Pd nanoalloys would greatly benefit from the present work by providing experimental data which can be used to refine of simulation methodology.

Fig. 1

 1 Fig. 1 (a) Low-magnification TEM image of carbon-supported Au-Pd nanoparticles (NPs) consisting of about 60 at.% Au. Prior to the transfer to the carbon support, the NPs were epitaxially grown, by pulse laser deposition, on a freshly cleaved NaCl(001) surface. In inset, the histogram of the size distribution of 386 NPs is shown. The average in-plane particle size is 4.8 ± 1.0 nm. In (b), the selected-area electron diffraction (SAED) pattern from the assembly of nanoparticles is presented. (c) and (d) UHRTEM images of two as-grown fcc NPs on the amorphous carbon support. The NPs are in close [001] orientation. The NPs are faceted single crystal with main edges parallel to the <110> directions and minor truncations in the <100> ones. The corresponding Fourier Transform (FT) patterns are also displayed alongside the corresponding UHRTEM images.

Fig. 2

 2 Fig. 2 (a) Low-magnification TEM images and the corresponding SAED patterns of two annealed specimen (samples A and B). Sample A consists of initially disordered fcc NPs, similar to those in Fig. 1, annealed at 600 • C during 10 hours. Sample B consists of NPs with composition comparable to those in Fig. 1 annealed at 500 • C during 8 hours. The size distribution of the two annealed samples are given in the insets of (a) and (b). The size distribution in (a) is constructed by analyzing the size of 240 NPs in sample A. The one shown in (b) is built from size distribution of 237 NPs from three different TEM images of sample B. (c) and (d) show SAED patterns obtained on samples A and B respectively. In (c), 001 and 110 superlattice reflections and a splitting of the 002 diffraction rings are visible.

Fig. 3

 3 Fig. 3 UHRTEM images of L1 0 (a-c) and L1 2 (d-f) Au-Pd nanoparticles from samples A and B. The corresponding FT patterns are shown below each image. (a-c) L1 0 -ordered NPs imaged close to the [100] (a) and [001] (b-c) directions. In those directions, chemical order is evidenced through the appearance of 001 and 110 superlattice reflections in the FT patterns. The L1 2 structures in (d-f) are in close [001], [101] and [112] orientations respectively. Again, strong superlattice reflections appear on the FT patterns.

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Acknowledgements

The authors acknowledge the support of the French Agence Nationale de la Recherche (ANR) under reference ANR-11-BS10-009. We are also grateful to Région Ile-de-France for convention SESAME E1845 for the support of the JEOL ARM 200F electron microscope installed at Paris Diderot University.