List of figures

List of tables

Introduction

With the rapid deve lopment of nanotechnology, colloida l semic onductor nanocrysta ls research is currently a topical area of scientific interest due to its excellent size dependent physical and optical properties. Nanocrystals are often referred to as "artificial atoms", which are characterized as adjustable emission wavelength according to the size, wide absorption range and fast luminescence decay time. Since the first controlled synthesis of colloidal nanocrystals in 1993 [START_REF] Murray | Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites[END_REF], the synthetic technique is continuous ly improved, making it poss ible to prepare the nanocrystals w ith various sizes, compos itions and shapes. In the past 20 years, breakthroughs have been made from the traditiona l II -VI and III-V groups to the very promis ing perovskite nanocrysta ls, and the synthesis of nanocrystals with different shapes, such as nanorods and nanoplatelets, are also enriching the application prospects. However, the increased Coulomb interaction between electrons and holes on the nanoscale significantly increase the non-radiative Auger recombination rate which becomes the ma jor pathway for charged exc iton a nd multiexc itons recombination, resulting in the decrease of fluorescence yield. This has become an important obstacle for nanocrystals in multi-carrier related applications. The appearance of heterostructures greatly improved the stability of nanocrystals, and also provided a great opportunity to s ignificantly reduce non-radiative Auger recombination. To date, nanocrysta ls have de monstrated the outsta nding optica l performances in a wide range of applications including TV screens, light-emitting diodes , lasing media and solar cells [START_REF] Schaller | High Efficiency Carrier Multiplication in PbSe Nanocrystals: Implications for Solar Energy Conversion[END_REF][START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF][START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Diroll | Violet-to-Blue Gain and Lasing from Colloida l CdS Nanoplatelets: Low-Threshold Stimulated Emission Despite Low Photoluminescence Quantum Yield[END_REF]. In addition, the potential applications of nanocrystals are still being explored, in which the scintillation materials in the form of nanocrystals have recently been received great attention in the scie ntific fie ld.

A scintillating material (scintillator) is a kind of converter transforming the energetic photon or partic le (ionizing ra diation) into low energy UV/vis ible photons. With the discovery of X-rays by Roentgen in 1895 [START_REF] Röntgen | On a New Kind of Rays[END_REF] and the emergence of the photomultiplier tube , the research of scintillators has made great progress, the scintillators have been discovered in various crysta lline and non-crystalline media, one of the most used of which is the thallium-activated sodium iodide (NaI:Tl). Up to date the scintillator-based radiation detectors are widely used in nuc lear me dic ine, high energy phys ics , homela nd security and industrial non-destructive testing [START_REF] Dujardin | Needs, Trends, and Advances in Inorganic Scintillators[END_REF]. The selection of the scintillation materia ls depe nds on the particular requirements of the cons idered application. With the deepening of the understanding of fundame nta l mechanis ms underlying sc intillation process, more and more materials are available and the properties of scintillators are constantly improving. At present, bulk materials still occupy the commercial scintillator market. In order to improve specific performances (timing, energy resolution...), new scintilla t in g materia ls are still a very hot research direction.

Research on radiation detection related to nanocrystals gradually emerged in recent years. The first configuration was reported in 2006 and applied to the detection of ga mma rays [START_REF] Létant | Semiconductor Quantum Dot Scintillation under γ-Ray Irradiation[END_REF]. Subsequent research on nano-scintillators focused on the combination of nanocrysta ls w ith bulk materials [START_REF] Hossu | X-ray luminescence of CdTe quantum dots in LaF 3 :Ce/CdTe nanocomposites[END_REF][START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF]. Recently, sc intillator in form of perovskite nanocrysta ls was demonstrated w ith remarkable performa nce [START_REF] Chen | All-inorganic perovskite nanocrystal scintilla tors[END_REF]. Unlike the traditional bulk scintillators, the output wave length of nanocrysta ls can be tuned spa n the vis ible light emission such that it can well match with the peak spectral sensitivity of the photomultiplier tube to improve the overall detection efficie ncy. Moreover, na noscintilla tors ha ve the advantages of liquid processing and convenient mass production, which can avoid the complicated procedures of high temperature fabrication and address the problem of high cost of scintillation materials. And some specific application benefit of the nanosize like in photodynamic therapy under X-rays [START_REF] Hu | Nanocomposite-Based Photodynamic Therapy Strategies for Deep[END_REF][START_REF] Kamkaew | Scintillat ing Nanoparticles as Energy Mediators for Enhanced Photodynamic Therapy[END_REF].

In spite of many attempts have been devoted so far on different types of nanoscintillators, the scintillation mechanism in nanocrystals is still not fully described. A few artic les have reported the process of energy deposition and relaxation in nanocrysta ls after energetic excitation. The interaction of materials with high-energy radiation basica lly produces a shower of high-energy electrons, which will multiply, thermalized and excite emitting centers. Due to the very small size, below the main migration distances of the involved spec ies, the energy depos ition and the distribution of exc ited states in na nocrysta ls w ill be significantly different from bulk materia ls. In addition, regarding the identification and quantification of excited states under high-energy electrons was reported [START_REF] Padilha | Response of Semiconductor Nanocrystals to Extremely Energetic Excitation[END_REF] , the preliminary observation indicated that after initial electron s hower, in the process of energy re laxation dow n to band sta tes within a nanocrystal, carrier multiplication and electrons scattering played an important role, leading to the formation of multiexcitons states and charged exciton states. Meanwhile, they als o pointed out that Auger recombination serious ly hindered the deve lopment of nano-scintillators. We believe that a more comprehensive understanding of the scintilla tion mechanis m in nanocrysta ls is needed to improve and discover nove l scintilla t io n materia ls.

Aim of the thesis

Motivated by potentia l of nanocrystals for radiation detection application, this thes is aims to explore the process of energy de pos ition and charging in nanocrysta ls under ionizing radiation with different II-VI heterostructures, inc luding spherica l quantum well and heterostructura l nanoplate lets. The excited states dynamics, especia lly Auger recombination and multie xc itons will be investigated w ith optica l spectroscopy techniques including steady state spectroscopy and time resolved technique. By comparing with intense optical excitation temporally and spectrally, we are able to verify and identify the formation of multiexc itons states under X-ra ys excitation. Meanwhile , through the investigation on the partic les with varying s izes, we w ill ana lyze the s ize dependence of the scintilla tion response, that is , the evolution of ratio of exc itonic spec ies with the particle size.

Organization of the thesis

The thesis is organized in 6 chapters as follow:

In Chapter 2, we present the fundamental concepts necessary to the understanding of this thesis. First, we introduce the quantum confinement effect and surface effect of colloidal semiconductor nanocrystals, and also introduce spherical quantum wells(SQWs), nanoplatelets(NPLs) and their heterostructures in detail. Secondly, we discuss the exciton dynamical processes in NCs, where mainly emphasize the formation and recombination of multiexc itons. Fina lly, we als o describe the sc intilla tion mechanis m of bulk scintilla tors and the development of nano-scintillators.

In Chapter 3, we first present the synthesis princ iples of core NCs and the heterostructures. Secondly, the characterization techniques used for the detailed investigation of the structura l and optica l properties in this thesis are briefly introduced. We also report the s ynthesis of CdS/CdSe/CdS SQWs w ith varying s hell thic kness, a nd different heterostructural NPLs, core/crown, core/shell and core/crown/shell, together with their characterization.

In Chapter 4, we first describe the steady-state radioluminescence spectroscopy, the time-resolved experiments. We also performed the spectral and dynamical measurements with various s ized SQWs under the optical and X-rays excitation. By comparing with simulation, we also presented the identification of the excitonic species during scintilla t io n process and further the investigatio n on their evolution with partic le size.

In Chapter 5, we describe the investigation of the scintillation performance of different II-VI nanoplatelets with steady-state spectroscopy and time-resolved methods, including core/crown, core/shell, and core/crown/shell NPLs in which three-sized core/shell NPLs was performed for the comparison.

The Chapter 6 concludes the work and discuss future outlooks. .

Chapter 2

Scientific background

In this chapter, before presenting our findings and related discuss ion, we first introduce the scientific background on colloida l semiconductor nanocrysta ls and the excited state dynamics , specially the Auger non-radiative recombination and the multiexc itons. Second, we briefly describe the scintillation mechanis m and the state of the art of the nanocrysta ls scintillators, a nd fina lly we introduce the optical and e lectronic properties of two kinds of nanocrysta ls : spherica l quantum wells(SQWs) and nanoplatelets(NPLs).

Colloid semiconductor nanocrystals

Colloida l semiconductor na nocrysta l(NCs) is a nano-s ized semic onductor materia l conta ining 100-100000 atoms. The size is smaller than or comparable to the exciton Bohr radius of the bulk material, therefore the NCs are subjected to quantum confinement effect, resulting in a quantized energy level and an adjustable energy gap. After first proposed by Ekimov in 1981 [START_REF] Ekimov | Quantum Size Effect in Three -Dimensional Microscopic Semiconductor Crystals ZhETF Pis[END_REF] , NCs have attracted great attention due to the ir fascinating chemical and phys ical properties, in whic h II-VI NCs are the most studied, such as CdSe, CdS and ZnS. The ability to precisely control the composition, size, and shape of II-VI NCs has been successfully realized and provides great flexibility in the engineering of the ir e lectronic a nd optica l properties. Also, by c ombining different materia ls , various types of heterostructures, like CdTe/CdSe , CdSe/ZnS, ha ve been fabricated to specifically alter the excitonic features of NCs to obtain desired characteristics. Nowadays, II-VI NCs has made outstanding progress in las ing media , single-photon s ources, photovolta ic convers ion, dis play devices, and biological imaging. With the continuous deve lopment of nanocrysta ls, its applications in other fields are constantly being explored.

Quantum confinement effect

An important parameter of a semiconductor materia l is the width of the energy gap that separates the conduction from the valence energy bands controlling the emiss ion luminescence wavelength [START_REF] Klimov | Nanocrystal Quantum Dots[END_REF]. When a semiconductor absorbs a photon with energy equal to or larger than its energy gap, an electron can be excited from the valence band to the conduction band, leaving a hole in the va lence band, and the formed e lectron-hole pair bound by Coulomb interaction is referred as an exciton. It has lowest energy state slightly be low the edge of the conduction band, which can recombine radiative ly to emit a photon [START_REF] Rabouw | Excited-State Dynamics in Colloida l Semiconductor Nanocrystals Top[END_REF]. For bulk semic onductors, the energy band structure depends on the ele ment compos ition and lattice structure and it is a fixed parameter. A natural length scale of the bounded electron-hole in the hydrogen like structure is given by the exciton Bohr radius , a B , which is determined by the strength of the electron-hole Coulomb interaction.

= * (2-1)
where ε is the dielectric constant of the material, is the Bohr radius of hydrogen atom

( =4 ℏ ⁄ ~ 0.53Å),
is the rest mass of the electron and m* is the reduced mass of the effective electron-hole pa ir. Exc iton Bohr radius is a materia ls intrins ic property. In this thes is , exc iton Bohr radii of semiconductors we used is 2. 9nm for CdS and 5.6 nm for CdSe.

In the case of semiconductor NCs with size comparable or smaller to the bulk exciton Bohr radius, confinement effect becomes important. Electronic energies depend directly on the degree of spatial confinement of electronic wave functions. One consequence of this is the vale nce and conduction band edges change from continuous to discrete energy leve ls, a nd the band ga p of the NCs gradua lly increases as the size decreases, as presented in Figure 2.1. This phenomenon is known as the quantum confine ment effect, which makes it poss ible to achieve continuous ly adjustable wavelength in NCs by simply tuning their size and shape [START_REF] Klimov | Nanocrystal Quantum Dots[END_REF][START_REF] Demir | Applied nanophotonics[END_REF]. The energy level structure of NCs can be calculated through the effective mass approximation theory, assuming a spherica l nanocrysta l w ith a radius of r, the bandga p of NCs can be expressed as [START_REF] Brus | Electronic wave functions in semiconductor clusters: experiment and theory[END_REF]:

= + ℏ + - . (2-2)
Here refers to the band gap of this materia l in bulk, ℏ is Planck's constant, m e and m h are the effective masses of electron and hole, respectively and e is the charge. is the dielectric constant in vacuum and is the relative die lectric constant of the material. The middle item on the right is the energy gap change caused by the quantum confinement effect, and the last item on the right is related to the electron-hole Coulomb interaction, which is the binding energy of exciton. The confinement energy is proportiona l to 1/r 2 whereas the exciton binding energy is proportiona l to 1/r leading to the dominance of the confinement effect. Therefore, we can see that within the confinement effect size range, as the size decreases, its energy gap will gradually become larger. Just by adjusting the size, nanocrystals with different energy gaps can be obtained.

On this bas is, the theoretical and experimenta l works have further studied the discrete atom-like energy levels of NCs and the corresponding transition selection rules [START_REF] Al | The Electronic Structure of Semiconductor Nanocrystals[END_REF]. The quantum confinement effect can occur in any dimension of a semiconductor materia l. D ifferent confinement dimens ions cause the different de ns ity of states of the semiconductor energy leve l. Figure 2.3 shows a schematic of the density of states as a function of dimens iona lity, going from three-dimens iona l (3D, bulk) to zerodimens iona l(0D, QDs) systems. In bulk materia ls , the charge carriers are free to move in all three dimens ions. The bands are square root dependent of the ir energy(~E 1/2 ). For a two-dimens iona l(2D) materia l, referred to as quantum well or nanoplate lets, only the thickness of this nanocrystal is sma ller than the exc iton Bohr radius and exc itons can freely move in the lateral directions. So the density of states is affected by the decreased number of freedoms for the charge carriers, resulting in step-like density of states. And for one-dimens iona l(1D) materia ls such as nanow ires and nanorods , the quantum confine ment effect plays an importa nt role in tw o dimens ions , this results in a saw-like dens ity of states with an inverse square energy dependence (~E -1/2 ). Fina lly, the semiconductor materia l confine d in all three spatia l dimens ions is referred to as quantum dots , they show discrete atom-like energy levels [START_REF] Rabouw | Excited-State Dynamics in Colloida l Semiconductor Nanocrystals Top[END_REF][START_REF] Edvinsson | Optical quantum confinement and photocatalytic properties in two-, one-and zero-dimensiona l nanostructures[END_REF]. This unique property caused by shape depende nce have greatly stimulated enormous interest in exploring nove l c lasses of materials with various electronic and optical characteristics.

Surface effect

Surface effect is another well-known nanoscale effect that causes different properties of NCs compared to bulk materials. The surface-to-volume ratio of NCs increases sharply as the particles size decreases. For bulk solids, their size is indeed much larger than the exciton Bohr radius of the particle, and most of the carriers are far away from the surface. Therefore, the influence of surface atoms on the material's properties is basically negligible for penetrating radiations, whereas the size of NCs is smaller than the excitons Bohr radius, carriers are easily affected by the surface. Therefore, NCs will be very sensitive to surface and environmental changes. The surface effect becomes more and more important on stability and optical properties of NCs with the further reduction of size. Compared with internal atoms, surface atoms have fewer neighbors, and therefore possess unsatisfied chemica l bonds , whic h is referred as dangling bonds. The increasing fraction of surface atoms leads to many dangling bonds, inducing defects on the surface. These defects will bring additiona l energy leve ls into forbidden band, cons idered as trap center to capture the charges created during an excitation, which is detrimental to the high quantum effic iency purs ue through non-radiative recombination pathway. In addition, the increase in dangling bonds will increase the surface free energy, leading to the instability, and NCs will tend to aggregate. This aggregation process is often irrevers ible , which greatly damages the optical properties of NCs and limits its application [START_REF] Donegá | Synthesis and properties of colloida l heteronanocrystals[END_REF][START_REF] Roduner | Size matters: why nanomaterials are different[END_REF]. Thus, reasonable elimination of dangling bonds and defects at the surfaces is often required to improve NCs performance.

To achieve the passivation of the surface, people utilize the ligands on the surface of NCs to form strong bonds with surface atoms to fill the dangling bonds, improving the chemical stability. The common binding groups of ligands are carboxyl, amino, and sulfhydryl groups. One end of the ligand is coordinated with the atoms on the surface, and the other end is compatible with the solvent to provide solution stability. The choice of hydrophilic and hydrophobic liga nds depends on the polarity of the solve nt. However, some types of liga nds ha ve been reported to quench optica l properties of NCs , it is important to choose the suitable ligands [START_REF] Munro | Quantitative Study of the Effects of Surface Ligand Concentration on CdSe Nanocrystal Photoluminescence[END_REF]. A lternative pass ivation method is to overcoat a shell composed of different materia ls , creating a heterostructure , which is referred as core/shell structure [START_REF] Dabbousi | CdSe)ZnS Core-Shell Quantum Dots: Synthesis and Characterization of a Size Serie s of Highly Luminescent Nanocrystallites[END_REF][START_REF] Reiss | Core/Shell Semiconductor Nanocrystals[END_REF]. The existence of a shell makes it possible to isolate the emiss ive core from its environment, which ma kes the NCs more stable and robust, and less sensitive to changes in its environment. The core/shell heterostructure will be discussed in detail in the followin g section.

Heterostructural nanocrystals

Heterostructural nanocrysta ls cons ist of a quantum dot core and a she ll of a (or more) distinct material. By using heterostructures instead of bare core NCs, the ability to create novel optoelectronic properties can be extended further. The core and the shell are typica lly c omposed of type II-VI and III-V semiconductors. This structure was demonstrated to address the low quantum yield issue of bare NCs with surface related defects by passivating the surface trap states. What's more, the formation of shell can isolate the emissive region with its environment, providing protection against environmental changes and photo-oxidative degradation [START_REF] Donegá | Synthesis and properties of colloida l heteronanocrystals[END_REF][START_REF] Reiss | Core/Shell Semiconductor Nanocrystals[END_REF].

Core/shell semic onductor nanocrysta l properties are based on the re lative conduction and va lence ba nd edge a lignme nt of the core and the she ll. The spatia l loca lization of the charge carriers in heterostructure can be manipulated by controlling the band offsets of the materials [START_REF] Reiss | Core/Shell Semiconductor Nanocrystals[END_REF]. Furthermore , the lattice parameters of she lling materia ls should be close to that of the core to reduce lattice stress created by the lattice mismatch and to facilitate epitaxial growth on the core. On the contrary, the growth of the shell results in strain, accompanied by the generation of defects at the core/shell interface or within the shell. These defects can act as trap states to capture the photo-ge nerated charge carriers and diminish the fluorescence quantum yield. As illustrated in Figure 2.4, according to the shell materia ls and the energy band alignment, core/s hell nanopartic les can be divided into three types:

(1) Type I: the conduction band energy of core material is lower than that of shell whereas the valence band energy of core is higher than that of shell. As a consequence, both electron and hole are confined within the core, making it poss ible to separate phys ica lly the surface of the optically active core from its surrounding medium. Due to the well surface passivation w ith respect to bare core, coating w ide band gap materials is us ua lly e mployed to reduce nonradiative recombination, which enhances photoluminescence quantum yield and the chemical stability against photo-degra dation. In addition to the improvement in quantum yield, shell growth is accompanied by a small red shift of the excitonic peak in the absorption and photoluminescence spectrum. This observation is attributed to a partia l leakage of the wavefunction into the shell material. The first published configuration was CdSe/ZnS [START_REF] Hines | Synthesis and Characterization of Strongly Luminescing ZnS-Capped CdSe Nanocrystals[END_REF]. However, the lattice parameter mismatch(12%) between CdSe and ZnS limit strongly the continuous growth of shell without the formation of interface defects [START_REF] Dabbousi | CdSe)ZnS Core-Shell Quantum Dots: Synthesis and Characterization of a Size Serie s of Highly Luminescent Nanocrystallites[END_REF].

(2) Quasi type-II: one carrier is delocalized over the whole nanocrystal, while the other is localized in one of the materia ls. The most investigated composition is CdSe/CdS. This allows the electron-hole spatial overlap to be tailored by controlling the size of shell. The delocalization of one carrier reduces the confinement effect, making it possible to tune the emission color over a large range. Therefore, the redshift in spectra upon she ll grow n is usua lly larger than that observed in type I. Moreover the lattice parameter mismatch of typical structure CdSe/CdS is only 4%, w hich conduces to the synthes is of the giant core/shell na noparticle [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF]. The reduction of e lectron-hole wavefunction overlap as the shell decreases the radiative transition probability and thus increases the exciton lifetime. Due to the decreased Coulomb interaction, Auger recombination rate is concluded to decrease in this configuration, making it feas ible for the investigation of multiexc ition properties and the suppression of blinking phenomenon [START_REF] Mahler | Towards non-blinking colloida l quantum dots[END_REF].

(3) Type II: In contrast to type I, both c onduction le vel and va lence leve l of one material are higher than the other one, therefore one of either the electron or holes is localized in the core while the other is confined in the shell. The separation of carriers leads to the s ignificant reduction in wavefunction overlap, like quasi type II, whic h has the effect on elongated lifetime of exciton. The emiss ion wave length is determined by the e nergy difference between the occupied states, which will be at a lower energy than either of the individua l energy ga p. This configuration thus has been de veloped in particular for near infrared emission.

The continuous de velopment of heterostructure enables it to satisfy multiple requirements of appropriate electronic and structural parameters simultaneously. For instance, multi-shell structure can realize the decrease of defected induced by lattice mismatch between the core and an outer shell by adding an intermediate layer, such as CdSe/ZnSe/ZnS [START_REF] Reiss | Low polydispers ity core/shell nanocrystals of CdSe/ZnSe and CdSe/ZnSe/ZnS type: preparation and optical studies Synth[END_REF]. Compared to the traditiona l core/she ll structure , it has proven to be more stable and higher quantum yie ld. Another structure is gradient core/she ll NCs with smooth interfacial potential [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF]. The further reduced steepness of the interfacial potential makes this configuration a great success in suppressing the Auger effect.

Excited state dynamics under photoexcitation

Figure2.5 is the s implified picture of the exciton dynamica l processes in NCs under photoexc itation, as mentioned above , the abs orption of a photon in nanocrysta ls can excite an electron from the valence band to the conduction band, leading to the formation of exciton. Subsequently, the exciton will undergo a series of processes, including relaxation and recombination, to dissipate the energy. This sequence of events involves time scales from a few femtoseconds to a few seconds after photoexcitation [START_REF] Rabouw | Excited-State Dynamics in Colloida l Semiconductor Nanocrystals Top[END_REF].

The poss ible pathways strongly depend on the specific exc itation conditions (excitation energy and intens ity). For the incident moderate energetic photon (E = ħꞷ > E g ), it will produce hot exciton with excess energy, typically the hot exciton rapidly loses its energy as heat or is trapped by defects. If the system is excited under high-energy photon (in princ iple E = ħꞷ ≥ 2E g ) or high-intensity excitation, the carrier multiplication process will take place, resulting in the formation of the multiexcitons, this will greatly complicate this whole dynamical process. Eventually, exciton is able to emit photon through radiative recombination or lose the energy via the non-radiative recombination, which is mainly caused by sample defects and Auger non-radiative recombination, relaxing the NCs back to the ground state. Here, the main possible excited state dynamics processes in NCs under different exc itation conditions are expla ine d in deta il in the follow ing sub-section. An understanding of the rates and pathways of exciton re laxation and recombination processes in NCs is very important for the comprehens ion of the fundamental physics as well as their utilizat io n in applications.

Hot carriers relaxation

When the excitation energy is equa l to the bandgap energy, it is referred to resonant excitation. On the other hand, in a general case, the photon energy is greater than the band gap, usua lly leading to both electron and hole occupy high leve ls in the conduction and deep ones in the va lence band, this process is known as non-resonant exc itation. A nd the carriers with excess energy are usually referred to as 'hot carriers'. In bulk semiconductors , the created hot carriers rapidly return to the band edge via phononassisted emission in an ultrafast process. This process is typically in hundreds of femtoseconds due to the strong coupling of carriers and phonons within the continuum of energy states. However, when we considered the nanocrystals, the cooling process is significantly different. As the size of the nanocrystals decreases, the discretization of electronic states appears near the band edges, and the electronic intraband spacing will continue to increase with the confinement till hundreds of meV, which is larger than the maximum phonon energy with only tens of meV, preventing the charge carrier relaxation [START_REF] Peterson | The Role of Ligands in Determining the Exciton Relaxation Dynamics in Semiconductor Quantum Dots[END_REF]. Auger-type energy transfer is generally accepted as the alternative process. Due to the larger effective mass of hole, valence band possesses a much higher density of state, making the smaller hole energy states spacing in valence band than electron energy states spacing in conduction band. Hence the electron transfers its excess energy to the hole through the Auger type energy transfer, accompanying with the relaxation to the edge of conduction band. Subsequently, holes diss ipate the energy rapidly with the assist of photons [START_REF] Nasilowski | Synthesis and optical spectroscopy of thick-shell semiconductor nanoparticles: applications to biological imaging[END_REF].

Exciton recombination

Once the electron and hole have relaxed to the ir respective band edge , subsequently excitons can recombine through radiative and non-radiative recombination. For radiative recombination, the exciton is released in the form of emitted photons. And the energy of emitted photon is equal to about the bandgap minus the bond energy. This characteristic rate is referred to as radiative rate . Sometimes the radiative recombination related to the trapped carriers may also occur, which is called trap emission. For non-radiative recombination, severa l types of this phe nomenon are poss ible in nanocrysta ls. Since the presence of defects and dangling bonds in nanocrysta ls , carriers are captured and blocked in these sites, the emiss ion can be quenched. In addition to the qualities of the partic les, Auger effect also is part of the origin of non-radiative recombination in charged and multiexcitons. The characteristic rate is referred to as non-radiative rate . Nonradiative pathways extre mely diminish fluorescence quantum yie ld. If we cons ider is the population of excited states at time t = 0, the evolution of this population will be defined by:

= - ( ) (2-3)
Here the rate k can be expressed by the fluorescence lifetime . The fluorescence lifetime is the time measured for the number of excited states to decay to 1/e of the original population via the loss of energy through radiative or nonradiative processes.

= + = + = (2-4)
So ( ) which is the number of excited states at time t is supposed in s imple case to be expressed by the single exponential decay equation of excited states as follow:

( ) = (0) = (0) (2-5)
Fluorescence quantum yie ld is one of the most important characteristics of fluorophores. It is defined as the ratio of the number of e mitted photons to the number of absorbed photons. Since the exciton recombination in nanocrysta ls us ua lly involves ma ny radiative and non-radiative processes, the quantum yie ld depends on the various recombination pathways and can be written as:

=

(2-6)

Auger recombination

Excitation of nanocrystals may lead to charged partic les and in such s ituation A uger recombination is the main non-radiative pathway for energy dissipation, which is mediated by Coulomb interactions between carriers. This process plays an important role in semiconductor nanocrystals when there are three or more charge carriers present, which are referred as trion and multiexcitons. Here, we take trion as an illustration, which is a lso calle d as charged exciton. The formation of charged exc iton is us ua lly attribute d to the charged carriers trapping or Auger ionization. A schematic diagra m Auger recombination in charged exciton system is illustrated in Figure 2.6: instead of emitting photons as neutral single exciton by radiative recombination, the energy is transferred to the additiona l carrier (electron in the conduction band), subseque ntly promoting it to higher excited energy leve l, and then diss ipated as heat by phonon-assisted thermal process. In direct-gap bulk semiconductors, this process needs to satisfy both energy and translational momentum conservation requirements and hence it exhibits a thermally activated behavior. Additionally, strong s patial overlap of wavefunctions for a ll carriers is required. Therefore, Auger recombination can only be observed at higher electron-hole pair dens ity. However, in nanocrysta ls, only energy-matching conditions are needed to be met because of the no well-defined trans lationa l momentum, together with the enhanced Coulomb interactions ow ing to the increased spatia l confine ment strongly improve the efficiency of Auger effect with respect to that of bulk materials [START_REF] Demir | Applied nanophotonics[END_REF], and it can occur for band edge carriers. The timescale of this process is typica lly on the order of 1-100ps , depending on exciton mult iplic ity, the size shape and compos ition, which is faster than that of radiative recombination. A general trend observed of Auger decay rate in core NCs is a linear depende nce of Auger lifetimes on NCs volume, w hich is referred as "universal volume scaling law" [START_REF] Klimov | Quantization of Multipartic le Auger Rates in Semiconductor Quantum Dots[END_REF]. There have been a few recent exceptions to this universal volume sca ling of Auger rates in 2D nanoplate lets [START_REF] Li | Thickness Dependent Bi-Exciton Auger Recombination in Colloida l CdSe Nanoplatelets: Breaking the "Universal Volume Scaling Law[END_REF]. The Auger effect in nanoplatelets was demonstrated more sensitive to the changes in vertical thickness(d) with a scaling factor of d 7 , indicating its strong depende nce on the degree of confinement. Meanw hile it increases linearly with its lateral area, reflecting the dependence of exciton-exciton collisio n frequency.

Basically, Auger process has always been considered as the dominant reason for the photoluminescence intermittenc y in s ingle quantum dots and the reduction of quantum yield [START_REF] Efros | Origin and control of blinking in quantum dots[END_REF]. Since the net result of this process is energy loss as heat, the Auger quenching is the main obstacle to commercial application of nanocrystals that require high efficiency, such as light-emitting diodes (LEDs) and lasing medium [START_REF] Rabouw | Excited-State Dynamics in Colloida l Semiconductor Nanocrystals Top[END_REF]. In the last decades, a lot of endeavors have been made to inhibit the Auger process. As discussed above , Auger lifetime is involved in spatia l confineme nt of charge carriers. Therefore, the Auger recombination rates can be reduced by increasing the size of nanocrystals w ith weak confinement complication. Besides, the degree of spatial overlap between the electron and hole wave functions is cons idered as another influenc ing factor. Heterostructures based on type II and quas i type II core/shell, have been continuous ly develope d to manipulate the overlap of electron and hole pa ir. Common configuration CdSe/CdS with thick she ll has showed outstanding performance on Auger effect suppress ion, the Auger rate can even be prolonge d to several nanoseconds [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF]. In addition, some research also report alloyed interface with smooth confinement potential leads to the suppressed Auger process, leading to a rather high quantum yie ld of mu lt ie xc ito ns [START_REF] Nasilowski | Gradient CdSe/CdS Quantum Dots with Room Temperature Biexciton Unity Quantum Yield[END_REF][START_REF] Park | Effect of the Core/Shell Interface on Auger Recombination Evaluated by Single-Quantum-Dot[END_REF]. Indeed, A uger recombination rate in nanocrysta ls strongly depends on the size, shape and potential steepness of the interface. Increasing requirements of high qua lity nanocrystals dema nds us to further understand the underlying phys ics of Auger recombination and reduce the Auger losses.

Multie xcitons recombination

As we discussed above , on the consequence of strong carrier-carrier Coulomb interaction in nanosca le, the efficie ncy of Auger non-radiative recombination is significantly enhanced, which represents the predominant pathway of multiexc iton decay [START_REF] Klimov | Quantization of Multipartic le Auger Rates in Semiconductor Quantum Dots[END_REF][START_REF] Klimov | Spectral and Dynamical Properties of Multiexcitons in Semiconductor Nanocrystals[END_REF][START_REF] Achermann | Multiexcitons confined within a sub-excitonic volume: Spectroscopic and dynamical signatures of neutral and charged biexcitons in ultrasmall semiconductor nanocrystals[END_REF]. The emergence of large-size and gradient NCs provided an excellent opportunity to exploitation of multiexc itons. The suppress ion of the Auger effect has greatly improved the fluorescence quantum yie lds of s ingle exc itons and multie xcitons of these NCs. And this breakthrough has provided a good opportunity to explore the huge potentia l of the multiexc itons phenome na in the application of semiconductor nanocrysta ls , such as low-threshold NCs lasers, light emitting diodes, as well as solarenergy convers ion [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF][START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF]. The realization of all of these applications requires detailed understanding of dynamica l and spectral properties of multiexc itons in nanocrystals. Hence, we introduce briefly the optical and electrical properties of biexciton and triexciton, and the carrier multiplication process.

Biexciton

The multiexcitonic state composed of two excitons in a single nanocrystal at the same time is the s implest configuration of neutra l multiexc itons , referred to biexc iton. Through nonradiative pathway, biexc iton annihilates an electron-hole pa ir and pass the energy to other carriers (either electron or hole), exciting the carriers to high energy state or even ionizing them out of the particles to form the charged exciton.

Multiplicity of recombination pathways accelerates the radiative lifetime of multiexc itons. The nonlinear scaling of radiative rate constants of N-exciton states ( ) with N was analyzed by K limov et a l. in 2008 [START_REF] Mcguire | N ew Aspects of Carrier Multiplication in Semiconductor Nanocrystals[END_REF]. They proposed "free-carrier" model and predicted experimentally and theoretica lly likely scales as N 2 in PbSe NCs. Therefore although the scaling factor of two has also been reported for biexciton [START_REF] Narvaez | Excitonic exchange effects on the radiative decay time of monoexcitons and biexcitons in quantum dots[END_REF], the relationship generally accepted between the exciton and biexciton radiative rates is τ 1r = 4τ 2r . On the basis of recombinations of the possible electron and hole angular momentum states, Katherine E. Shulenberger theoretically and experimentally explained the origin of statistica l scaling factor of four. D ue to angular momentum selection rules, only half of the excitons are available to radiatively recombine in exciton, and in biexciton two spinpaired exc iton can recombine independently, giving rise to a radiative speed up by a factor of two compared to a single spin-paired exciton. The combination of the factors results in a biexciton radiative rate four times faster than that for the exciton [START_REF] Shulenberger | Multiexciton Lifetimes Reveal Triexciton Emission Pathway in CdSe Nanocrystals[END_REF]. Recently, the deve lopment of the suppression of Auger effect makes the fluorescence quantum yield of single excitons and biexcitons reach nearly 100% [START_REF] Nasilowski | Gradient CdSe/CdS Quantum Dots with Room Temperature Biexciton Unity Quantum Yield[END_REF][START_REF] Park | Near-Unity Quantum Yields of Biexciton Emission from CdSe / CdS Nanocrystals Measured Using Single -Particle Spectroscopy[END_REF]. The lifetime of bie xciton also composes with radiative and A uger non-radiative items, = + . Therefore, according to the statistical scaling, the emission quantum yield of biexc iton is expressed in Eq 2-7, allowing us to calculate it directly with the exponential fitting parameters of the decay curves.

= = = = (2-7)
To further understand the spectral and te mpora l properties of biexc iton, the interaction between exc iton and exc iton is a nother important problem to investigate, which is often define d as Δ BX = E BX -2E X , where Δ BX is interaction energy, and E X and E BX are single exciton and biexciton energies, respectively [START_REF] Piryatinski | Effect of Quantum and Dielectric Confinement on the Exciton-Exciton Interaction Energy in Type II Core/Shell Semiconductor Nanocrystals[END_REF]. The magnitude and s ign of the interaction energy depe nd on the spatia l distribution of electrons and holes , which is the electrons and holes wavefunction overlap integra l. Therefore, different energy ba nd alignments will lead to different Coulomb interaction energies. Specifically, in type I nanocrysta ls , both electrons and holes are restricted in the same area, as shown in Figure 2.7(a), the interaction force between excitons is attraction, which corresponds to the negative binding energy of biexc itons , leading to a red shift w ith respect to s ingle exc iton on the spectrum. It has been observed in bare nanocrysta ls and type I samples such as CdSe/ZnS [START_REF] Achermann | Multiexcitons confined within a sub-excitonic volume: Spectroscopic and dynamical signatures of neutral and charged biexcitons in ultrasmall semiconductor nanocrystals[END_REF][START_REF] Caruge | Transient photoluminescence and simultaneous amplified spontaneous emission from multiexc iton states in CdSe quantum dots[END_REF]. On the contrary, in the quasi-type II and type II structures, the alignment of the energy bands causes the difference in the distribution of electrons and holes, which leads to the imbalance of local charges. The repulsive force between excitons and exc itons has been reported in CdSe /CdS, CdS/ZnSe and other structures, as seen in Figure 2. 7(b), and a blue shift re lative to s ingle exc iton is caused [START_REF] Htoon | Highly Emissive Multiexcitons in Steady-State Photoluminescence of Individua l "Giant" CdSe/CdS Core/Shell Nanocrystals[END_REF]. However, in few recent studies, they achieved to control the nature and strength of exciton-exciton interactions by adjusting the core-shell sizes with CdSe/CdS structure. The biexciton shift transiting from red-shift to blue-shift regimes in a single material system has been observed, typically, thin shell was demonstrated to exhibit the attractive interaction due to the negligible delocalization of the electron wave function. Therefore, the tunable amplified spontaneous emission(ASE) has been realized by engineering biexciton interactions [START_REF] Cihan | Attractive versus Repulsive Excitonic Interactions of Colloida l Quantum Dots Control Blue-to Red-Shifting (and Non-shifting) Amplified Spontaneous Emission[END_REF][START_REF] Gollner | Random Lasing with Systematic Threshold Behavior in Films of CdSe/CdS Core/Thic k-Shell Colloida l Quantum Dots[END_REF]. Taking direct carrier-carrier Coulomb coupling and interface polarization into account, Andre i Piryatinski [START_REF] Piryatinski | Effect of Quantum and Dielectric Confinement on the Exciton-Exciton Interaction Energy in Type II Core/Shell Semiconductor Nanocrystals[END_REF] theoretically calculated the biexciton binding energy of type II structure by perturbation the ory. Due to the large degree of spatia l separation between electron and hole , the repuls ion energy in type II can reach to the order of 100meV. While in quas i-type II, because of the delocalization of one carrier, the relative ly enlarged wavefunction overlap of electrons and holes reduces the interaction energy, on the order of 10 meV.

Trie xciton

In comparison with biexciton, the situation of triexcitons is extremely complicated, because it involves not only the lowest energy quantized state, but also a higher P energy leve l. According to Pauli's exclus ion princ iple , the conduction ba nd of CdSe is composed of the 5s orbita l of cadmium, so the lowest quantized state 1Se at the bottom of the conduction band can only conta in tw o electrons w ith oppos ite spins. While the top of valance band is ma de up of six degenerate 4p atomic orbita l of sele nium, ignoring the spin-orbit c oupling, the lowest energy leve l 1S 3/2 state of the hole is four-fold dege neracy and can accommodate up to four holes. Since the effective mass of the hole is larger, the energy le ve l in terva l of t he va le nce ban d is sma lle r, p-like hole states can be populated as a result of thermal activation in the triexciton ground state [START_REF] Achermann | Multiexcitons confined within a sub-excitonic volume: Spectroscopic and dynamical signatures of neutral and charged biexcitons in ultrasmall semiconductor nanocrystals[END_REF][START_REF] Franceschetti | Radiative Recombination of Triexcitons in CdSe Colloida l Quantum Dots[END_REF]. The occupation of the 1P state can only be stable when the 1S state is fully filled. Based on these studies, P-level related emission will appear at high-energy positions with respect to single excitons, approximately to 200meV [START_REF] Fisher | Room-Temperature Ordered Photon Emission from Multiexciton States in Single CdSe Core -Shell Nanocrystals[END_REF].

Similar to biexciton, the efficient Auger process seriously hinders its application. In the past 20 years, thanks to the deve lopment of ultrafast optica l instruments , such as timeresolved spectroscopy, high-e nergy triexc iton emiss ive feature has been wide ly reported [START_REF] Caruge | Transient photoluminescence and simultaneous amplified spontaneous emission from multiexc iton states in CdSe quantum dots[END_REF][START_REF] Fisher | Room-Temperature Ordered Photon Emission from Multiexciton States in Single CdSe Core -Shell Nanocrystals[END_REF]. What's more, on the basis of the study regarding Auger effect suppress ion, Florenc io García-Santamaría et al. observed the pronounced change of emiss ion color in amplifie d sponta neous emiss ion experime nts with giant CdSe/CdS NCs [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF]. And the yield of white light emiss ion was observed by Miche l Nasilowski et.al, suggesting the effective impr ove ment of high-order(>2) multie xcitons efficienc y by synthesized gradient interfacial NCs [START_REF] Nasilowski | Gradient CdSe/CdS Quantum Dots with Room Temperature Biexciton Unity Quantum Yield[END_REF]. ground-state biexciton, (c) or via the nominally forbidden 1Pe-1Sh, also followed by a thermal process. Finally, each of these three pathways can bring the system to the ground-state biexciton.

Most of the current st udies on triexc itons are based on the differences in spectra, and the predominant recombination pathways for trie xcitons rema in poorly understood. The generally accepted decay mode is a cascade stepped decay: TX-> BX-> X-> ground state. But the specific process from TX->BX is still very contr overs ia l. As shown in Figure 2.8, there are three possible recombination pathways, which are S-like: 1S e -1S h , and P-like: 1P e -1P h , 1P e -1S h . Most reports so far acquiesce that this process is completed by P-like transition [START_REF] Fisher | Room-Temperature Ordered Photon Emission from Multiexciton States in Single CdSe Core -Shell Nanocrystals[END_REF][START_REF] Fisher | Multiexciton fluorescence from semiconductor nanocrystals[END_REF]. In 2018, Katherine E. Shulenberger [START_REF] Shulenberger | Multiexciton Lifetimes Reveal Triexciton Emission Pathway in CdSe Nanocrystals[END_REF] demonstrated that S-like and Plike occur simultaneously, and the S-like trans ition is dominant through the third order correlation of single CdSe/CdS NCs. In the same year, Jaehoon Lim [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF] discovered that the 1P lasing threshold is slightly lower than the theoretical value for the 1P e -1P h transition. They attributed it to the occurrence and dominance of the 1P e -1S h transition pathway. According to the transition selection rule, 1P-1S trans ition is prohibited. However, due to close proximity between the corresponding energy leve ls in the va lence band, the energy difference between 1S 3/2 and 1P 3/2 is a lmost ne g lig ib le , even on a la rge size scale of nanocrystals, the effect of mixing of the band-edge 1S and 1P hole states can be significant, leading to the permiss ion of 1P-1S transition. In addition, this asymmetric transition has als o been prove n in many other reports regarding the charged biexc iton a nd trie xc it on [START_REF] Achermann | Multiexcitons confined within a sub-excitonic volume: Spectroscopic and dynamical signatures of neutral and charged biexcitons in ultrasmall semiconductor nanocrystals[END_REF][START_REF] Htoon | Highly Emissive Multiexcitons in Steady-State Photoluminescence of Individua l "Giant" CdSe/CdS Core/Shell Nanocrystals[END_REF]. However, in the spectral profile, these two different P-like transitions in CdSe nanocrystals are difficult to be distinguished.

Carrier multiplication

Generally, absorbing one photon w ith moderate energy(E g <E<2E g ) create only one hot exciton, the excess energy will be released as heat by interacting with phonons(Figure 2.9(a)). When a single high-e nergy photon(>2E g ) is absorbed, carrier multiplication(CM) has been observe d in bulk semiconductors s ince the 1950s and expla ined by impact ionization [START_REF] Vavilov | On photo-ionization by fast electrons in germanium and silic on[END_REF], where the hot electron dissipates the energy by pumping the carrier cross the bandgap, forming the additional electron-hole pair, as illustrated in Figure 2.9(b). However, in bulk semic onductors , d ue to strict momentum conservation and very effective phonon-assisted attenuation, the efficiency of carrier multiplication is considerably low. The no well-defined trans lationa l momentum and the strong Coulomb interaction between carriers make the carrier multiplication process in NCs could be highly efficient [START_REF] Klimov | Spectral and Dynamical Properties of Multiexcitons in Semiconductor Nanocrystals[END_REF]. Since the efficient carrier multiplication effect is of great significance for improving the photoe lectric convers ion effic iency of solar cells , photodetectors , etc., the CM in nanocrystals has received extensive attention. For traditional solar cells, since the excess energy is not effective ly used, the maximum theoretical photoe lectric convers ion efficiency of a single-gap cell is about 33.7%, whic h is the Shockle y-Queisser limit [START_REF] Shockley | Detailed balance limit of efficiency of p-n junction solar cells[END_REF]. However, considering the carrier multiplication effect, the highest photoelectric convers ion effic iency is predicte d to be as high as 44% [START_REF] Klimov | Detailed-balance power conversion limits of nanocrystalquantum-dot solar cells in the presence of carrier multiplication[END_REF]. In 2004, the V.I.Klimov research group first discovered the efficient carrier multiplication effect in PbSe nanocrysta ls us ing trans ient absorption spectroscopy [START_REF] Schaller | High Efficiency Carrier Multiplication in PbSe Nanocrystals: Implications for Solar Energy Conversion[END_REF]. Subsequently, highly effic ient carrier multiplication effects have also been observed in a variety of nanocrystals, such as CdSe and perovskite CsPbI 3 NCs [START_REF] Hu | Carrier Multiplication in a Single Semiconductor Nanocrystal[END_REF][START_REF] De Weerd | Efficient carrier multiplication in CsPbI3 perovskite nanocrystals[END_REF]. Multiexcitons created by carrier multiplication is the major content we investigate in this subject. Therefore, the brief introduction here will contribute to the understanding of following discussion section.

Scintillation process and scintillator 2.4.1 Scintilla t ion process

X-ray is a kind of electromagnetic radiation w ith very short wave length between ultraviolet and gamma rays, about (20 0. 06)×10 -8 cm. It was discovered by the German physicist W.K. Roentgen in 1895, so it is also called Roentgen Ray [START_REF] Röntgen | On a New Kind of Rays[END_REF]. The device that generates X-rays is called an X-ray tube, which c ons ists of three parts: cathode, anode and window, as shown in Figure 2.10. The two electrodes are used for the target materia l (such as W, Ag, Mo, Cu) bombarde d by electrons and the fila ment (generally W wire) that emits electrons. Both electrodes are sealed in a high-vacuum glass or ceramic hous ing. The power supply part of the X-ra ys tube inc ludes a lowvoltage power supply for heating the filament and a high-voltage generator for applying high voltage to the two electrodes. When the fila ment passes enough current to generate an electron cloud, and there is enough voltage (kV) to be a pplie d between the anode and the cathode, the electrons will hit the target materia l at high speed. When the high-speed electrons reach the target surface, the movement is suddenly blocked. The lost kinetic energy will be released in the form of photons to form a continuous part of the X-rays spectrum. The radiation generated in this form is called Bremsstrahlung. In addition, by increasing the acceleration voltage, the energy carried by the electrons increases, and it is poss ible to e ject the inner electrons of the anode meta l atoms. As a result, holes are formed in the inner layer, followed by the outer layer electrons transition back to the inner layer to fill the holes, and accompanied by emitting photons with a wavelength of about 0.1 nm. Since the energy released by the outer electron transition is quantize d, it forms a characteristic line in the X-rays spectrum, which is called Characteristic radiation. By adjusting the fila ment c urrent, the te mperature of the filame nt and the amount of electrons emitted can be changed, thereby changing the tube current and the X-rays photon flux. D ifferent excitation pote ntia ls or different target materia ls can produce X-rays of different energies. Due to the bombardment of high-energy electrons, the temperature of the X-rays tube is very high during operation and forced cooling of the anode target is required. The approaches to detect X-rays and more generally ionizing radiation are divided into direct detection and indirect detection [START_REF] Jung | Flexible Gd 2 O 2 S:Tb scintilla tors pixela ted with polyethylene microstructures for digital x-ray image sensors[END_REF]. In direct detection X-rays are converted to electrica l signa ls directly. Although this method is highly accurate, in many situations, these detectors are not easily applicable due to their cost and the required cryogenic application. So the indirect method is now more widely used. Xrays are first converted into a visible light, and this conversion is called as scintillation, which is followed by the conversion of light signa ls into electrical s igna ls in photodetectors, as shown in Figure2.11. Scintillators are thus the materials that can convert high energy radiation such as X or gamma-rays into a visible light, and they play a very important role in the field of radiation detection, such as medica l imaging, homela nd security, high energy phys ics calorimetry and so on [START_REF] Dujardin | Needs, Trends, and Advances in Inorganic Scintillators[END_REF][START_REF] Glodo | New Developments in Scintilla tors for Secur ity Applications[END_REF]. A sc intillation detector essentia lly is composed of a scintilla tor materia l and a photodetector that can convert the outcoming light to an electrical signal. The scintillators can be generally categorized into organic scintillators and inorganic sc intillators : organic sc intillators c omprise such as organic liquids a nd plastic sc intillators, and inorganic scintillators typica lly cons ist of s ingle crysta ls , powder or thin films of large band-gap sa lts matrix dope d w ith activator atoms of appropriate energy states, such as NaI:Tl, CsI:Tl. Nowadays, inorganic scintillators are more widely used in numerous fie lds , and the nove l inorga nic scintillators w ith superb qua lities are continuously investigated to satisfy the various applications.

The very first inorganic scintillator material NaI:Tl, was discovered in the late 1940s [START_REF] Hofstadter | Alkali Halide Scintillation Counters[END_REF]. Despite the emergence of considerable new scintillator materials after that, NaI:Tl is still the first choice for many applications where cheaper, larger size, and higher efficiency scintillator is needed. However, it suffers from the common shortcomings of inorganic scintillator, long decay constants and hygroscopicity due to limitations in their intrinsic material properties and fabrication requirements. Since that time, a diversity of scintilla tors has been reported for various applications. For instance, CsI:Tl is not hygroscopic w ith high light yie ld [START_REF] Van Sciver | Scintillations in Thallium-Act ivated Ca I 2 and CsI[END_REF], but it shows the poor spectral match with PMT and extremely long decay time. On the other hand, Bi 4 Ge 3 O 12 (BGO) with higher stopping power, making it a very efficient absorber has been widely used in high-e nergy phys ics applications [START_REF] Weber | Luminescence of Bi 4 Ge 3 O 12 : Spectral and decay properties[END_REF] , but the light yield is much lower than that of NaI:Tl and CsI:Tl. The relative ly new materia ls based on Ce 3+ doping and Pr 3+ doping possess the fast decay time (typically tens of ns) and high quantum effic iency [START_REF] Van Loef | High-energy-resolut io n scintilla tor : Ce3+ activated LaBr3[END_REF]. But the high price due to raw materials and processing severely restricts their widespread usage.

Overall, each material possesses unique properties, the choice of scintillator materia ls depends on the practica l application require ments. The most important characteristic s considered for a scintilla tor sele ction are as follow [START_REF] Nikl | Scintilla t ion detectors for x-rays[END_REF]:

-stopping power: determines how efficiently a material will absorb the radiation during the convers ion stage of sc intillation. H igh stopping power means large fraction of the incident rays and partic les can be depos ited in the materia ls. As mentioned a bove, this parameter is strongly re lated to the dens ity and effective atomic number of the scintillator, especially for the photoelectric interaction. Therefore, materials with high density and high Z are desirable. -light yield: corresponds to the number of photons emitte d per unit of depos ited energy in the material. It is expressed in photons/MeV, is given by the relation: LY = 10 6 •SQ/(βE g ), where E g is the bandgap, S is the efficiency of electron-hole transport to the optical center, Q is the luminescence quantum efficiency of the optical center and β is a constant. H igh light yie ld is very beneficia l to reduce the noise in medica l imaging.

-response time: corresponds to the time of emitting a photon after radiation absorption. Short decay time is very important for applications where fast timing is required such as time-of-flight applications.

-radioluminescence spectrum: it should well match the spectral sensitivity of the photodetectors. -energy resolution: the ability of the scintillator to distinguish between different radiation energies. It is very important for the spectral measurement.

In addition, large scale production, chemical and radiation stability, self-absorption and cost are also key factors that determine the choice of a particular scintillator. Actually, the selection of scintillator is according to specific requirements and application fie lds. No scintilla tor can satisfy all the requirements for use.

The physical process of scintillation is a rather complex issue in inorganic scintilla tor, whic h can be divided into three consecutive sub-processes [START_REF] Nikl | Scintilla t ion detectors for x-rays[END_REF]: (I) conversion, (II) transport and (III) luminescence, as depicted in Figure 2.12.

In the first step, high energy photons interact primarily with materials through ionization events, followed by the creation of a large number of highly energetic electrons and holes formed in the material. The interaction of an electromagnetic radiation w ith the matter occurs as three types depending upon the incident photon energy of the incident radiation in scintillation: photoelectric effect and Compton scattering are dominant mechanisms for low (be low a few hundred keV) to medium energies (a few hundred keV-1MeV) respectively, and the interaction of the radiation with the matter is governed by the electron-pos itron pa ir creation for energies above 1.02 MeV [START_REF] Maddalena | Inorganic, Organic, and Perovskite Halides with Nanotechnology for High-Light Yield X-and γ-ray Scintillators Crystals[END_REF].In thi s thesis, we will focus mainly on detection of photons in the X-rays, involving the photoelectric effect and Compton scattering. The linear absorption coefficient of the photoelectric absorption for energy scales approximately with Z 4 . Hence, increasing the atomic number will increase the absorption probability exponentially, which is consistent with the need of heavy atoms in scintillator [START_REF] Dujardin C | [END_REF]. Afterwards, the hot electrons and deep holes that are created in this process will undergo the multiplication process through electron-electron inelastic scattering and Auger processes. When their energy is lower than the threshold, all the electrons and holes will be thermalized to the lowest states in the energy band.

In the transport process, electrons and holes in the conduction and va lence bands are transported towards the luminescence centers, respective ly. The undes ired impurities and lattice defects during the fabrication of material will induce lots of traps in the forbidden gap. These traps compete to capture the migrating charge carriers, causing the cons iderable loss of excitons through the non-radiative recombination. In addition, the long delay is also the result of these intermediate states. After trapping, a certain thermal energy enables release of the charge from the shallow traps, and then these released carriers will continue to participate in the radiative recombination, resulting in considerable delay or afterglow. Optimization of fabrication and preparation procedures can significantly decrease the defects and traps within a material and improve the scintilla t in g qualitie s .

In the last step, the electrons and holes are captured by the luminescence centers, and then undergo a recombination process, giving rise to a desired luminescence light. In some particular groups of materia ls the lig ht generation occurs in radiative trans itions between the valence and first core bands [START_REF] Nikl | Scintilla t ion detectors for x-rays[END_REF].

Nanocrystals scintilla tor

In recent years, the increasing demand for radiation detection materials in many applications has led to extensive research on scintillators. The discovery of new scintillators has always been an important subject. Nanocrystals have received extensive attention due to their excellent optical properties [START_REF] Dujardin | Needs, Trends, and Advances in Inorganic Scintillators[END_REF]. Unlike the traditional bulk scintilla tors , the output wavelength of na nocrysta ls can be tuned to span the vis ible light emission such that it can well match with the peak spectral sensitivity of the photomultiplier tube to improve the overa ll detection effic iency. Furthermore, the solution-processability and practical scalability can also reduce the manufacturing cost and difficulty, ma king it a promis ing materia l for radiation detection. In the past two decades, the response of a large number of nanocrystals under high-energy electrons, alpha particles, X-rays and gamma rays has been widely studied.

In 2006, S.E. Letant and T.F. Wang for the first time reported the characterization of CdSe/ZnS core/shell NCs in porous glass matrix under α particle and ga mma-rays radiations. They demonstrated that the scintillation efficiency of the CdSe/ZnS NCs is two times better than that of the conventional NaI:Tl scintillator [START_REF] Létant | Semiconductor Quantum Dot Scintillation under γ-Ray Irradiation[END_REF][START_REF] Létant | Study of porous glass doped with quantum dots or laser dyes under alpha irradiation[END_REF]. This strongly suggests that higher effic ienc ies and resolutions are possible with nanocrysta l-based scintilla tors , and motivated the further research on the scintillator materia ls in the form of nanocrystals.

The relatively low stopping power is one of the shortcomings of II-VI nanocrystals , leading to a weak scintillation luminescence, hence the combination of a matrix with high stopping power and na nocrysta ls has been usua lly used to investigate the ir potentia l to be a new class of scintillator. Several research have been reported to fabricate NCs embedde d nanocompos ite scintillators, the improving of e nergy efficie ncy has been achieved via energy transfer mechanisms between matrix and NCs. The chosen NCs materia ls should well match the emiss ion spectral region of the matrix. The high stop ping power of the matrix can compensate low energy absorption of NCs. Mingzhen Yao and Marius Hossu integrated CdTe NCs with LaF3:Ce nanoparticles and observed the enhanced NCs fluorescence upon both UV and X-rays excitation [START_REF] Hossu | X-ray luminescence of CdTe quantum dots in LaF 3 :Ce/CdTe nanocomposites[END_REF][START_REF] Yao | Luminescence enhancement of CdTe nanostructures in LaF3:Ce/CdTe nanocomposites[END_REF]. In addition, the NCs-polymer was also fabricated to investigate the scintillation behavior. CdTe NCs was embedde d into polyvinyl a lcohol matrix, together with the 2,5-diphenyl oxazole (PPO) as exciton harvesting [START_REF] Thacker | Quantum Dot-Organic Polymer Composite Materials for Radiation Detection and Imaging[END_REF], and the addition of CdS NCs in polyvinyl toluene (PVT)-PPO polymer enhanced scintillation from a 1.03 times to more than 3.6 compared to us ing PVT-PPO without any NCs [START_REF] Tam | Quantum-dot doped polymeric scintilla t io n material for radia tion detection[END_REF]. However, the intensity of the scintillation light scales with the NCs concentration. High concentration usually induces strong reabsorption and the scattering due to agglomeration. And this type of NCs-polymer hybrids were also designed for ga mma-ray sens or, s ignificant enhancement in photoc onductivity induced by the gamma radiation in nanocomposite has been demonstrated [START_REF] Hosseini | Dosimetric investigation of a new quantum dots/nanocomposite (CdTe QDs/PVK) sensor for real-time gamma radiation detection[END_REF].

Environmenta lly friendly nanocrystals are always the deve lopme nt trend of nanotechnology, to be sc intilla tor, it w ill be benefic ial to decrease the damage to human tissues, facilitating the practical application in medical imaging. ZnS NCs doped with magnesium sulfide, ZnS(MgS) was studied by E. R. Burke, showing good proportionality and high sensitivity of X-rays [START_REF] Burke E R, Dehaven | Scintillat in g quantum dots for imaging x-rays (SQDIX) for aircraft inspection 42ND ANNUAL[END_REF]. In 2018, Eder Jose Guidelli et al. synthesized ZnSe NCs using a microfluidic reactor with uniform sizes and excellent optical properties. They reported the radioluminescence from ZnSe cores under X-rays exc itation, revea ling the possibility to be a low-toxic scintillator [START_REF] Guidelli | Mechanistic Insights and Controlled Synthesis of Radioluminescent ZnSe Quantum Dots Using a Microfluidic Reactor[END_REF]. However, due to the low mass density and effective atomic number of Zn-based NCs, their absorption of high-energy rays is very low, whic h severely limits the scintilla ting performance.

Photo-instability is one of the key issues of semiconductor nanocrystals, which severely degrade the performance of scintilla tion. Several studies reported the optica l response of NCs to high energy photons, a rapid loss of output light was observed when the NCs was exposed under X-rays or γ-rays [START_REF] Withers | Rapid degradation of CdSe⁄ZnS colloidal quantum dots exposed to gamma irradiation[END_REF][START_REF] Gaur | Influence of Ionizing Radiation and the Role of Thiol Ligands on the Reversible Photodarkening of CdTe/CdS Quantum Dots[END_REF]. Afterwards, Girija Gaur et al. reported the significantly reduced emission intensity with CdTe/CdS QDs immobilized within porous s ilica under ionizing irradiation [START_REF] Gaur | Influence of Ionizing Radiation and the Role of Thiol Ligands on the Reversible Photodarkening of CdTe/CdS Quantum Dots[END_REF]. They also pointed that photoionization, carrier traps, and Auger recombination play a role in the radiationinduced photodarkening. Through a thiol-rich surface re-passivation, near complete recovery of X-rays induced emission intens it y was achieved, but only partia l luminescence recovery was observed for gamma-irrad ia te d samp le s, wh ic h is the resu lt of permanent reduction caused by lattice displacements in the NCs. In 2016, Marie-Eve Delage et al. demonstrated multishell QDs are more resistant to ionizing radiation than core/shell QDs, especially in the kV energy range, further confirming the importance of the surface passivation [START_REF] Delage | Robust shell passivation of CdSe colloidal quantum dots to stabilize radioluminescence emission[END_REF].

In addition to the traditional spherica l quantum dots with different compos itions , nanocrystals of different shapes have also received extensive attention. The first successful synthesis of a two-dimens iona l II-VI CdSe nanoplatelets(NPLs) was in 2008 [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF]. Due to its fascinating properties, it has recently been explored in the field of radiation detection [START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF][START_REF] Turtos | Ultrafast emission from colloida l nanocrystals under pulsed X-ray excitation[END_REF]. In 2019, Turtos R. M. et al. constructed hybrid NCs-bulk scintillators structure to measure the scintillation response for 511 keV gamma -rays [START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF]. A thin film of CdSe/CdS core/crown NPLs was deposite d on standard LY SO bulk scintilla tor, in w hich NPLs is in charge of fast time response while the bulk sc intilla tor provides stopping power and scintillation light. Ultrafast decay of 80ps assoc iated to multiexc itons has been e xtracted as the pre liminary result, providing a strong impetus for future in-depth research.

The nove l nanostructure in the form of perovs kite nanocrystals has emerged to be a promising scintillation material in recent years. It shows remarkable performance as scintilla tor benefit from high stopping power and defect tolerance for high quantum yield [START_REF] Chen | All-inorganic perovskite nanocrystal scintilla tors[END_REF][START_REF] Zhang | Metal Halide Perovskite Nanosheet for X-ray High-Resolution Scintilla t ion Imaging Screens[END_REF]. In 2018, Chen et al. prove d 3D perovs kite CsPbBr 3 na nocrysta ls has high fluorescence efficiency under X-rays excitation. The emission wavelength can be easily tuned across the visible light by tailoring the anionic component. In addition, the intense radioluminescence decreases significantly the detectable limitation of the X-rays dose, reducing the damage to human tissues caused by radiation [START_REF] Chen | All-inorganic perovskite nanocrystal scintilla tors[END_REF]. Furthermore , the solutionprocessability and practical scalability make them suitable for the mass scale production. However, the inherent thermal and chemical instability and the heavy metal toxicity are the obstacle for this type peroskite nanocrystals in practical applications.

Although a variety of nanomaterials are active in the field of radiation detection, the process of energy deposition and relaxation in semiconductor nanocrystals is still ambiguous and few reported s o far. As discussed above, the inc idence of high-energy photons yields a predictable cascade of energetic electrons, since the absence of transport process in direct gap nanocrysta ls semic onductor, the hot electrons and holes are directly therma lized towards the conduction and va lence band edges. Unlike bulk materials , nanocrysta ls be have as discrete energy leve ls related to quantum c onfinement, making the excitonic species unknown after the hot carriers ultimately relax to the edge of the energy band.

For the energy depos ition, in 2010, Anne-Laure Bulin et al. used Monte Carlo simulations to numerically estimate the spatial energy distribution resulting from the interaction between the gamma-rays photons and the Gd 2 O 3 nanocrysta ls. They demonstrated that a s ignificant fraction of the energy is depos ited out of the particle , even for the diameter larger than 100nm [START_REF] Bulin | Modelling energy deposition in nanoscintillators to predict the effi-ciency of the X-rayinduced photodynamic effect 9[END_REF]. To date, there is no direct study on unexamined steps in energy relaxation down to energy band states within a NCs excited by X-rays after the initial electron shower. Instead, the Lazaro A. Padilha research group [START_REF] Padilha | Response of Semiconductor Nanocrystals to Extremely Energetic Excitation[END_REF] has compared the temporally and spectrally resolved radiative recombination produced follow ing photoexc itation w ith 20keV energetic electrons ionizing radiation, identifying and quantifying the types of e xcitation states produced in NCs under high e nergy excitation. The analysis directly reveals a dominant population of charged-and multi-excitonic states, consistent with large contributions from carrier multiplication, multiple types of scattering and ionization processes during relaxation. They a lso pointed out that the Auger effect is a key factor restricting the study of excited states under ioniz ing radiation. The investigation of cathodoluminescence can provide guidance for the exploration of X-rays-induced exc iton recombination. In our work, tha nks to the development of c hemica l synthes is techniques, the reduction of Auger effect gives us a good opportunity to realize the fine observation and identification of multie xcitons states under X-rays excitation in nanocrystals.

Spherical quantum wells and nanoplatelets 2.5.1 Spherical quantum wells introduction

Despite the apparent promise of thic k-she ll NCs on both Auger recombination suppress ion and enhanced absorption cross-section, however, the thickness of the shell is not allowed to increase indefinitely. The intrinsic reason is the difference in lattice constants between the two materials of the core and shell, namely, with the increase in the thickness, the structural stress on the interface will continue to increase. When the thickness of the shell reaches a critical value, the further growth will introduce structural defects into the nanostructures, and these defects will capture carriers like surface traps, resulting in a serious decrease in the fluorescence efficiency of the NCs. In the commonly used synthetic materials of core/shell nanocrystals, the lattice parameter mismatch between CdSe and ZnS reaches 12%, such a huge amount of misfits make them not suitable choices even for traditional thick-shell NCs. The better candidates for the synthesis of thick-shell NCs should be CdSe and CdS because of the well energy band alignment and the only 4% lattice mismatch between two materials. However, according to the calculation of the critical thickness, when the CdSe layer grown on top of the CdS core exceeds around 3. 6nm [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF], the quantum yie ld w ill drastically reduce due to the formation of misfit defects.

In order to solve this limitation, spherical quantum wells(SQWs) are considered as an alternative structure. It combines the excellent characteristics from both nanoplatelets and regular core/s hell nanostructures. The first SQWs synthes ized through s uccessive ion layer absorption and reaction (SILAR) was reported by Peng's group, ultrathin layer of semiconductors w ith a narrow ba ndgap, e pitaxia lly grown onto a nanocrystal template with a substantia lly w ider bandgap [START_REF] Battaglia | Colloidal Two-Dimensional Systems: CdSe Quantum Shells and Wells[END_REF]. Since the emissive layer is thinner than the critical thickness, it enables the alleviation of structural stress between two materials by strained layer epitaxial growth. SQWs are considered as an optimal model to create large nanoparticles w ith a less defect growth. In addition, the thick she ll layer ensures the stability and reduces the probability of the SQWs being charged by reducing the possibility of carriers being trapped by the surface traps. Therefore, it could achieve the significantly enhancement of quantum yield compared to the traditional NCs. In 2016, Jeong et al. synthes ized CdS/CdSe/CdS SQWs through dropw ise precursor method with near unity quantum yie ld [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF]. On the contrary to the NCs , the quantum yie ld of SQWs show a continuous enhancement with increasing shell thickness. As the best investigated configuration, CdS/CdSe/CdS SQWs comprise three layers : a thin CdSe layer sandw iched between a CdS core and several la yers of CdS overcoating , as depicted in Figure 2.13(a). The electronic properties exhibit the strongly core, well and shell thickness dependence. Spec ifically, SQWs were also reporte d to possess a quasitype-II band alignment (Figure 2.13(b)), in which a hole is strongly localized within the CdSe well, whereas an electron is delocalized over the entire nanostructure volume. The electron confinement energy reduces with increasing exterior CdS shell thickness, resulting in the red shift of the abs orption a nd the PL peak pos ition [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF], as shown in Figure 2.14(a). And the same tendency of energy shift was observed for different core and well sized SQWs [START_REF] Xu | Exciton radiative recombination in spherical CdS⁄CdSe⁄CdS quantum-well nanostructures[END_REF]. Meanwhile, the decrease in the electron-hole overlap integra l with increasing CdS she ll thickness results in a progress ive reduction of the s ingleexciton recombination rate( In 2020, Gabrie l Nagamine et al. presented that SQWs enable a relatively large absorption cross-section w ithout the quantum yie ld loss , the quantum yie ld of SQWs about 40% larger than counterpart QDs with the same absorption cross-section [START_REF] Nagamine | Efficient Optical Gain in Spherical Quantum Wells Enabled by Engineering Biexciton Interactions[END_REF]. Since absorption cross-section is proportiona l to the volume of nanocrysta ls, the predicted large absorption cross section in SQWs is attributed to that its volume could far exceed that of traditiona l QDs. This promis ing property is benefic ial for tremendous research application, achieving the low threshold amplified spontaneous emiss ion and lasing, reducing the working current when applie d to e lectroluminescent devices and reducing the damage to the tissue when they are used in biological imaging [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Nagamine | Efficient Optical Gain in Spherical Quantum Wells Enabled by Engineering Biexciton Interactions[END_REF].

In addition, another remarkable property of SQWs is the suppression of A uger nonradiative recombination, leading to the long multiexcitons lifetime. Analogous to QDs, biexc iton lifetime depends strongly on the shell dimens ion, which can be attributed to the increase in the delocalization of the electron and hole wave functions. Till now, the biexc iton lifetime of SQWs was demonstrated on the order of nanosecond [START_REF] Nagamine | Efficient Optical Gain in Spherical Quantum Wells Enabled by Engineering Biexciton Interactions[END_REF][START_REF] Kholmicheva | Sustained Biexciton Populations in Nanoshell Quantum Dots[END_REF]. Meanwhile, as mentioned above, biexciton binding energy is determined by excitonexciton Coulomb interaction. Through modifying the ratio of different regions , one can manipulate the s ign and magnitude of bie xciton binding energy in SQWs , the trans ition from attractive to repulsive force was observed [START_REF] Nagamine | Efficient Optical Gain in Spherical Quantum Wells Enabled by Engineering Biexciton Interactions[END_REF].

Owing to the thick she ll, SQWs usually exhibit outsta nding optical and chemica l stability. By applying silica encapsulation via reverse micro-emulsion method, the stability of SQWs can be further improved, and it can be dissolved in polar solvents, which is convenie nt for bio-labe ling or displa y applications [START_REF] Wang | Highly luminescent silic a -coated CdS/CdSe/CdS nanoparticles with strong chemical robustness and excellent thermal stability[END_REF]. To date, due to the large absorption cross section and the Auger effect suppression, the research of SQWs mostly focused on the optical gain media and light emitting diodes [START_REF] Nagamine | Efficient Optical Gain in Spherical Quantum Wells Enabled by Engineering Biexciton Interactions[END_REF][START_REF] Porotnikov | Low -threshold la ser medium utiliz in g semiconductor nanoshell quantum dots[END_REF]. In this thesis, motivated by the fascinating properties, we will utilize this structure to investigate the scintilla tion behavior of nanocrysta ls on a large-scale (up to 40nm in dia meter) under Xrays excitation, this will give an ins ight to the exciton recombination mechanis m under extremely energetic excitation.

Nanoplatelets introduction

As mentioned earlier, with the development of nanocrystalline synthesis methods, nanomateria ls of different structures have been realized. In 2008, S. Ithurria successfully synthesized two-dimensional CdSe nanoplatelets for the first time [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF]. Subsequently, nanoplatelets with different composition have also been successfully fabricated [START_REF] Diroll | Violet-to-Blue Gain and Lasing from Colloida l CdS Nanoplatelets: Low-Threshold Stimulated Emission Despite Low Photoluminescence Quantum Yield[END_REF][START_REF] Pedetti | Optimized Synthesis of CdTe Nanoplatelets and Photoresponse of CdTe Nanoplatelets Films[END_REF]. Their outsta nding optica l and e lectronic properties have attracted widespread attention. Different than other classes of NCs, generally these colloida l NPLs feature lateral dimens ions on the order of tens of nanometers with the atomic -scale vertical thickness, while the 2D Bohr radius of the CdSe NPLs is cons idered to be around 3.5nm [START_REF] Tessier | Spectroscopy of Single CdSe Nanoplatelets[END_REF] , leading to the strong quantum confinement effect in only one dimens ion. By changing the synthesis reaction parameters, atomica lly precise control of the vertica l thic kness can be achieved.

The width of the fluorescence spectrum is composed of two factors: homogeneous broadening and inhomogeneous broadening. The coupling of photons and phonons will lead to homogeneous broadening of the spectral lines, but the inhomogeneous broadening is determined by the s ize distribution of the synthes ized nanocrysta ls. High monodispers ity of NPLs a llows to e liminate the inhomogeneous broadening, and they thus exhibit pure ly homogenous broadening. This results in a very sharp PL profile w ith a full-width-at-half-maximum (FWHM) typically narrower than 10 nm. In addition, NPLs possess gia nt osc illator strengths resultin g from increased exciton center of mass coherent motion [START_REF] Ithurria | Colloida l nanoplatelets with two-dimensional electronic structure[END_REF]. This huge oscillator strength enables significantly enhanced absorption cross-section and ultrafast recombination rate.

Controlling the number of monolayers (ML) significantly affects the absorption and photoluminescence spectra of the NPLs. Due to the emission wavelength of CdSe NPLs is only determine d by the thickness, it genera lly has fixe d emiss ion wavelengths . Figure 2.15 shows the spectra of several CdSe NPLs with different thicknesses, 3 ML, 4 ML and 5 ML, which have been commonly investigated for almost a decade, emitting at around 463 nm, 513nm, and 553 nm. The quantum confine ment is relaxed w ith the increasing thickness, causing a red shift of the excitonic features. For 4ML NPLs, three typical features are observed on the absorption spectrum, they correspond to the heavyhole(~512nm), light-hole(~480nm) and split off transitions (~420nm). The excitonic features remain a lmost in the same spectral pos ition regardless of the lateral dimens ions because they are much larger than the exciton Bohr radius. When the thickness of the CdSe NPLs is as thin as 3ML, the surface defects will lead to intensive emission. The strong surface related deep trap photoluminescence emission is more obvious in thinner NPLs, which is attributed to the unpassivated atoms at the surface [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF]. However, bare core NPLs usually suffer from low quantum yield and poor stability, which seriously affect the practical application of NPLs. Therefore, in order to increase quantum yield, improve optical and chemical stability, and achieve a wider spectral coverage, a divers ity of NPLs-based heterostructures have been produced in the past decade, for example core/crown, core/shell and core/crown/shell as illustrated in Figure 2.16, making these two-dimens iona l materials have made great progress in ma ny applications.

The most widely studied CdSe/CdS core/crown heterostructure has been synthesized by M. D. Tessier and coworkers in 2014 [START_REF] Tessier | Efficient Exciton Concentrators Built from Colloida l Core/Crown CdSe/CdS Semiconductor Nanoplatelets[END_REF]. The CdS crown region was grown only in the lateral direction without changing the vertical thicknesses of the NPLs. Since the quantum confine ment of the core is retained, the growth of the crown does not affect the excitonic features of the NPLs. The growth of laterally extended crown is manifested as an absorption peak at around 405 nm, leading to a strongly increase in the absorption cross section. Ow ing to the lateral passivation, the core/crown has been reporte d w ith the significant improvement of quantum yield. In addition, the core/crown structure NPLs with different compositions has also received extensive attention, such as type II CdSe/CdTe, alloyed CdSe/CdSeTe [START_REF] Gao | Low-threshold lasing from colloida l CdSe/CdSeTe core/alloyed-crown type-II heteronanoplatelets[END_REF][START_REF] Kelestemur | Type-II Colloida l Quantum Wells: CdSe/CdTe Core/Crown Heteronanoplatelets[END_REF].

In addition to core/crown, inspired by core/shell QDs structure, the low quantum yield and poor stability of NPLs are expected to be improved by coating a shell on the core NPLs. The growth of shell in the vertical thicknesses will induce the relaxation of the quantum confinement effect, making the red shift of excitonic features. Therefore, by controlling the shell thickness, the optical properties can be easily manipulated. Meanwhile, due to the enhanced exciton-phonon coupling w ith these materia l combinations, the excitonic features are slightly broadened with respect to core-only NPLs, as shown in Figure 2. 16(b). In recent years, different composition core/shell NPLs has been fabricated by different synthetic methods [START_REF] Altintas | Highly Stable, Near-Unity Efficiency Atomically Flat Semiconductor Nanocrystals of CdSe/ZnS Hetero-Nanoplatelets Enabled by ZnS-Shell Hot-Injection[END_REF][START_REF] Rossinelli | Compositiona l Grading for Efficient and Narrowband Emission in CdSe-Based Core/Shell Nanoplatelets[END_REF], in which a lloys and/or gra dient complex heterostructures have been demonstrated the near-unity quantum yield. Moreove r, ow ing to the partia l separation of e lectron and hole wavefunctions and the smooth confinement potentia l, Auger recombination has been significantly suppressed in this hererostructure.

In addition, more complex core/crown/shell structures have also been employed in the several researches [START_REF] Kelestemur | Platelet-in-Box Colloidal Quantum Wells: CdSe/CdS@CdS Core/Crown@Shell Heteronanoplatelets[END_REF][START_REF] Shendre | Ultrahigh-efficiency aqueous flat nanocrystals of CdSe/CdS@Cd 1-x Zn x S colloida l core/crown@alloyed-shell quantum wells[END_REF]. They pointed out that when compared to core/shell NPLs, core/crown/shell NPLs exhibit significantly enhanced absorption cross-section in the blue spectral region ow ing to crown layer and improved brightness due to better passivation of the sidewalls of the inner core. So far, the hot injection method for this architecture has still not been reported. The synthesis based on c-ALD technique usually causes low QY, only Shendre et al. presented ultrahigh-efficiency aqueous colloidal core/ crown/alloyedshell quantum wells with QY up to 90% in water [START_REF] Shendre | Ultrahigh-efficiency aqueous flat nanocrystals of CdSe/CdS@Cd 1-x Zn x S colloida l core/crown@alloyed-shell quantum wells[END_REF]. Even though NPLs can be synthesized by we ll-controlle d synthes is procedures and their excellent properties have received considerable attention, there are still controversies about the growth mechanism of the two-dimensional structure. Since the first synthes is of CdSe was due to the addition of short-chain salts in the reaction solution [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF] , short-cha in ligands were regarded as the key factor in the latera l growth of nanocrysta ls from the beginning. Long-cha in and short-cha in ligands (such as oleate and acetate salts) passivated the large and small cadmium-rich surfaces of NPLs, respectively. The long-chain ligand stabilizes the large facet, while the short-cha in ligand makes the small facet more unstable, allowing the NPLs to grow laterally and maintain a constant thickness. Lately, Riedinger et al. synthes ize d CdSe NPLs in me lts w ithout the presence of short-cha in carboxylate liga nds , they pointe d out that the addition of short-chain carboxylates is critical only in standard liqu id-phase protoc ols, leading to the reduction of monomer concentration, and the solubility of the monomer affects the shape of the crystals. Meanwhile, they proposed a theory, regarding the growth mechanism of zinc blend CdSe NPLs, based on an intrinsic instability in growth kinetics leading to highly anisotropic shapes [START_REF] Riedinger | An intrinsic growth instabilit y in isotropic materials leads to quasi-two-dimens iona l nanoplatelets[END_REF]. The nucleation barriers on narrow facets are smaller than on wide facets, therefore, the growth rates on lateral and vertical dimensions are orders of magnitude difference, yie lding the anisotropic-shaped NPLs even for materials with is otropic crystal structures. On the other hand, Peng's group [START_REF] Chen | Symmetry-Breaking for Formation of Rectangular CdSe Tw o-D imens iona l Nanocrysta ls in Zinc-Blende Structure[END_REF] has suggested a different mechanism based on the oriented attachment of seeds through intraparticle ripening.

Chapter 3 Colloidal Synthesis and Characterization of Spherical Quantum Wells and Nanoplatelets

The fascinating properties of nanocrystals have received worldwide attention. Thanks to the continuous deve lopment of the colloid synthetic approach, nanocrysta ls of different sizes, shapes, and compos itions have been successfully synthes ize d in the past decade[101] , and the optica l and e lectronic properties of na nocrysta ls have been significantly improved. Meanwhile, in comparison with bare core, the emergence of heterostructures with higher stability and quantum yield provides better opportunities for their applications in various fields [START_REF] Donegá | Synthesis and properties of colloida l heteronanocrystals[END_REF][START_REF] Reiss | Core/Shell Semiconductor Nanocrystals[END_REF].

Therefore, this chapter is mainly divided into four parts, we will first present the synthesis theory and setup for core and the core/shell structures. Secondly, the optical and structural characterization techniques used in this thesis are briefly introduced. The last two parts are about the specific synthes is protocols and characterization of nanocrystals , including spherica l quantum wells(SQWs) and different heterostructura l na noplatelets (NPLs), core/crown, core/shell and core/crown/shell.

Synthesis theory and setup

Nanocrystals nucleation and growth

Since the emergence of qua ntum dots , many efforts have been put in the understanding of the formation mechanis ms of mono-disperse nanocrysta ls. This progress ive comprehens ion makes it poss ible to achieve syntheses of nanocrystals with controlla ble size , shape, and compos ition, in which the therma l decompos ition is the most successful nanocrysta ls preparations in terms of qua lity and monodis persity [START_REF] Klimov | Nanocrystal Quantum Dots[END_REF]. The involved nucleation and growth are presented in the following sub-section.

Gibbs -Thompson theory

Homogeneous nuc leation is a phase change process that is the spontaneous formation of a new thermodyna mic phase from a meta-stability state [102]. Thus , in contrast to heterogeneous nuc leation, the occurrence of homogene ous nuc leation needs to overcome higher energy barriers. The thermodynamic theory of formation of the new phase is classical Gibbs-Thompson theory, as shown in Figure 3.1.

The Gibbs free energy of a nuc leus is in genera l expressed as the sum of a bulk term ΔG 1 that is proportional to the volume of the nucleus, and a surface term ΔG 2 , that is proportional to its surface area We a s sume that the nucleus is spherical in shape with radius r, hence, total Gibbs free energy is given as equation 3-1,

∆G = ∆G 1 + ∆G 2 = 4 3 πr 3 •∆G v + 4πr 2 •γ (3-1)
where γ is the surface energy per unit area and ΔG v is the free energy per unit volume in the bulk. The volume term is negative whic h is the driving force for nuc leus formation.

And the surface term has a positive contribution due to the creation of a solid and liquid interface. From the total Gibbs energy curve, we can see that the minimum radius of formed nuc leus is r c , wh ic h is referred as crit ic al s ize. The nuc le us s ma lle r than r c will dissolve into the solution, and the nucleus larger than r c can exist for further growth. The critical size is reached when free energy ΔG gets to a maximum, which is:

d∆G/dr=0 (3-2)
Simplifying the above result, we can derive that

r c = - 2γ ∆G v (3-3) ∆G c = 16πγ/3(∆G v ) 2 (3-4) 3.1.1.

Lamer burst nucleation theory

Burst nucleation theory was proposed by Lamer and co-workers in 1950s[103]. Origina lly Lamer theory described the production of mono-dispersed sulfur brines. The model is used today for growth of most diverse nanopartic les. It is based on homogeneous nuc leation theory, and separates the nucleation process and growth process. Nucle i are generated at the same time and subsequently grow w ithout addit iona l nucleation. Thus, the crystallization of the nanoparticles is complete and the particle size distribution is more uniform. However, if we do not stop the reaction at a reasonable time, the further growth of nanoparticles occurs in stage IV due to Ostwa ld ripening, in which large partic les compete w ith sma ller ones. Large nanopartic les are thermodynamica lly favora ble, a nd they will be continuous ly grow ing ow ing to the diss olution of sma ller partic les [104]. Hence, poly-dispersed samples are often attributed to Ostwald ripening process.

According to the discuss ion above , reaction factors have significant influence on the qua lity of synthesized samples. First, the temperature pla ys a key role during the synthesis, high temperature is required to form the nuclei, however the size and shape of nanocrystals are less controllable in this case, and temperature also determines the reactivity of the precursors. Therefore, a suitable temperature is extremely important for the precise control of nanocrystals. Second, if the precursor is introduced into the reaction by injection, the rate of injection is another key parameter to adjust the synthesis. The size distribution of nanocrystals can be controlled by fast or s low injection. Fast injection introduces the abrupt increase of the monomer to reach a very high super-saturation, leading to a narrow size distribution of nanocrysta ls in a very short period of time [103]. In addition, an appropriate annealing time after growth stage enables minimize the influence of Ostwald ripening on size dispersion.

Nanocrystals synthesis setup and method

The chemical synthesis reaction usually takes place in a three-neck round-bottom flask as illustrated in Figure 3.3. One neck of the flask is connected with a reflux condenser to limit the evaporation of the solvent during heating, one with a rubber septum through w hich precurs or could be injected or products re moved, and the third with a temperature probe. The flask is ma inta ined under an inert gas atmosphere (argon) and heated by a heating mantle. All the chemicals in the flask need to be degassed in order to achieve water and oxygen free conditions. And stirring during the reaction can make precursors mixed efficiently and avoid the aggregation of the nanocrystals. As mentione d above, the mono-dis perse nanocrysta ls s ynthetic strategies have been investigated s ince quantum dots were proposed[105,106]. To date, the thermal decomposition is the most effective and successful method for high quality nanocrystals synthesis. This method can be divided into two techniques : hot-injection and noninjection. The hot-injection method was introduced by Bawendi and co-workers on the synthesis of cadmium cha lcoge nide na nocrysta ls [START_REF] Murray | Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites[END_REF]. Sharp burst nucleation occurs due to the rapid injection of precurs or into hot solution, inc luding solvent a nd s urfactants. Thus, hot-injection enables the production of narrower s ize dis persion nanocrysta ls. However, it is difficult to scale up the synthes is in large quantities. Alternative ly, the deve lopment of non-injection method addressed this proble m, in which the precursors, solvent and surfactants are mixed in the flask at a low temperature, and then heated up to the nuc leation temperature. In order to achieve the formation of mono-s ized nanocrystals , refraining from the conc urrent nucleation and growth, the precursor choice and the setting of reaction temperature are extremely important during this gradually raising temperature process [START_REF] Yang | Synthesis of CdSe and CdTe Nanocrystals without Precursor Injection Angew[END_REF] 3.1.2 Core/shell structure synthesis Core/shell nanocrystals research has been one of the hot topics in recent decades. On the one hand, the core/shell structure can effectively eliminate the dangling defects on the surface of the bare core, reduce non-radiative transitions, and improve luminous performance. On the other hand, the core/shell structure composed of semiconductor nanocrysta ls w ith different properties has new functions and important application prospects in many fields.

The core/shell structure was synthes ized for the first time in 1996, Guyot-Sionnest et al. successfully covered a ZnS shell on the core CdSe, achieving the enhancement of quantum yield and stability [START_REF] Hines | Synthesis and Characterization of Strongly Luminescing ZnS-Capped CdSe Nanocrystals[END_REF]. Since then, two approaches were developed to synthesize core/shell structures. The SILAR (successive ion layer absorption and reaction) method based on the formation of one monola yer at a time by alternately injecting the anion and cation shell precursors. This method successfully solved the proble m of se lf-nuc leation of the she ll precursor. In 2003. via SILAR, Peng's research group successfully s ynthes ized up to 5 monolayers CdS shell on 3.5nm core CdSe with narrow PL line w idths in the range of 23-26nm [START_REF] Li | Large-Scale Synthesis of Nearly Monodisperse CdSe/CdS Core/Shell Nanocrystals Using Air-Stable Reagents via Successive Ion Layer Adsorption and Reaction[END_REF]. However, the SILAR method is a very time-consuming procedure, which is not conducive to the formation of thick shells.

Alternative ly, the other method is dropw ise, that is continuous injection of she ll precursors at the same time into the flask containing the core. This requires a good control of the temperature and the slow rate of injection to avoid secondary nucleation of shell material. If the injection rate is too slow, it will lead to Ostwald ripening of the core. The amount of precursors can also be calculated to predict the final size of the nanocrysta ls. In comparis on to the SILAR method, the continuous injection is much faster, but usually results in relatively less monodisperse samples [START_REF] Reiss | Highly Luminescent CdSe/ZnSe Core/Shell Nanocrystals of Low Size Dispersion[END_REF]. In our thesis, the aim is to prepare the particle with ultra-thick she ll. In order to s imulta neous ly satisfies the criter ia : time -saving and high uniformity, we utilize the continuous injection method and meanwhile using the precursors with relatively low reaction rate to achieve a wellmaintained particle size distribution during shell growth.

Optical and structural characterization techniques

To study the rationality of the synthetic method, different characterization techniques are needed to collect sufficient information on the synthesized nanocrystals, such as the size, shape, composition and the excitonic features. The investigation of the qua lity of the synthes ize d nanocrysta ls can give a guide line to improve the synthes is and the deve lopme nt of a new protocol. In this thes is , the structural techniques used are transmission electron microscopy (TEM), and the optical techniques include absorption spectroscopy(ABS), photoluminescence spectroscopy (PL), photoluminescence excitation spectroscopy (PLE).

Absorption spectrum

Absorption spectroscopy is a powerful technique to determine the absorption behavior of the synthes ized nanocrys tals , reflecting the size and compos ition of the sample. In an absorption spectrum, quantum confinement effect is well manifested as the red shift of the excitonic peaks as nanocrystal size increases, which corresponds to the decreasing energy band, as shown in Figure 3.4. And then the bulk be havior a ppears in large particles. Therefore, we can determine nanocrystals s ize from the pos ition of first excitonic peak [START_REF] Yu | Experimental Determination of the Extinction Coefficient of CdTe, CdSe, and CdS Nanocrystals[END_REF]. In addition, the different exc itonic peaks in s pectrum revea l the optica l transitions between the discrete energy levels due to quantum confine ment effect. The main absorption peak (1S 3/2 (h)->1S 1/2 (e)) exhibits a large oscillator strength, leading to a clear feature. And less pronounced absorption features in the shorter wavelength range correspond to exc ited states of higher energy, where the energy leve ls behave as a quas i-continuum [112]. The width of these peaks is strongly affected by the size distribution of the nanocrysta ls dis persion, polydis perse samples induce an inhomogeneous broadening of the excitonic features.

Photoluminescence spectrum

Photoluminescence (PL) is the property of a material to emit light after excitation by absorbing photons. In semic onductors , abs orbing photons w ith an energy superior to its band gap can excite the electron to the conduction ba nd, leading to the formation of exciton. The radiative recombination of exciton leads to the light emission. The emitted photons have an energy corresponding to the band gap of the semiconductor. Therefore, PL spectrum is usually utilized to follow the evolution of semiconductor nanocrystals growth during the synthesis. Both dispers ion solution sample and s olid film sample can be used in this measurement. Similar to absorption spectrum, the size-dependent emiss ion color is associated with quantum confinement effect. As illustrated in Figure 3.5, by changing the size of CdS/CdSe quantum dots, the emission spectra can be tuned to span the whole visible light range. Furthermore, the position of the emission peak corresponds to the band-edge energy, and the width of the PL spectrum is characterized by FWHM. The FWHM of the peak can also give us the information about the size distribution of samples [START_REF] Rogach | Semiconductor nanocrystal quantum dots: synthesis, assembly, spectroscopy, and applications[END_REF]. It should be noted here that the defects on II-VI nanocrystals surface usua lly induce a trap energy leve l in the forbidden band, therefore trapped carriers emission take place on the longer wavelength than band edge emission. Proper surface passivation is necessary to get rid of these trap states, otherwise the fluorescence spectra can be strongly impacted by the trap emission, as depicted in Figure 3.5.

Typically, the emission peak is red-shifted compared to the corresponding excitonic feature in the absorption spectrum. This phe nomenon is us ua lly referred to as Stokesshift, and this shift originates from the fine structure of the exciton energy levels. The self-absorption caused by a small Stokes-shift is one of the reasons for poor performance on lasing media, scintillator and others application where high concentrations of nanocrystals are needed [START_REF] Tam | Quantum-dot doped polymeric scintilla t io n material for radia tion detection[END_REF][START_REF] Taghipour | Sub-single exciton optical gain threshold in colloidal semiconductor quantum wells with gradient alloy shelling[END_REF]. Therefore, an enlarged Stoke-shift can be required to reduce self-absorption effect. And for the experiments, in order to avoid the selfabsorption, typically, the absorbance should be below 0.1 for the spectrum range we measured.

Photoluminescence excitation spectrum

Photoluminescence excitation spectroscopy(PLE) is a complementary technique to identify the origin of PL peak feature. At a selected emission wavelength, the intensity of photoluminescence is measured at varying excitation wavelengths. We could observe all the excitonic states contribution to the luminescence of the selected wavelength. It is expected that PLE spectra of the nanocrystals with well-defined size, shape and compos ition should be similar to their absorption spectra, as shown in Figure 3.6(a). Therefore, comparing to absorption spectrum makes it possible to figure out the presence of a secondary population. If the sample is polydis perse and secondary population appears during the synthes is process, a huge difference can be observed between ABS and PLE spectra (Figure 3.6( b)). The solution samples should be dilute d to avoid the selfabsorption, otherwise the PLE value will be much lower than absorbance. Usually, it is better to keep the absorbance lower than 0.1 at the excitation wavelength range.

In addition, the PLE measurements taken at different e miss ion wavelengths provides valuable information about the uniformity of the synthesized nanocrysta ls. Generally, the detected wavelengths are fixed at the lower energy emission side, the higher energy emiss ion s ide and the peak of the emiss ion, as shown in Figure 3.6(c). When the three PLE spectra show exactly similar excitonic transitions, it suggests the formation of highly monodispersed and uniform na nocrysta ls and the effect of inhomogeneous broa de ning on linewidth is negligib le .

Quantum yield measurement

Fluorescence quantum yie ld (QY) is the emiss ion effic ienc y of nanocrysta ls, whic h the ratio of the number of e mitted photons by abs orbed photons , as expressed in e quation (3-5): Q = Photons em / Photons abs [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF][START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Diroll | Violet-to-Blue Gain and Lasing from Colloida l CdS Nanoplatelets: Low-Threshold Stimulated Emission Despite Low Photoluminescence Quantum Yield[END_REF] where Q is the quantum yie ld. Two approaches are adopted for measurement of quantum yield, relative and absolute methods.

-Relative method

This method is w idespread to determine the nanocrysta ls solution quantum yie ld. It is performed under identical measurement condition to compare an unknown sample with a known qua ntum yie ld fluorophore , which is cons idered as a reference [116]. At the same absorbance, to compare the emission intensity with standard is a reliable way to estimate the quantum yield. Much more precise calculation is the gradient method, as depicted below:

Q = Q ref •( Grad Grad ref )•( n 2 n ref 2 ) (3-6)
where Grad is the gradient obtained from the plot of the integrated fluorescence intensity vs. optica l dens ity, the subscript ref denotes reference, and is the refractive index of the solvent.

The absorption and emission region of the reference is better if similar to test sample. To a vo id se lf-abs orption effects, for 10 mm cuvette, the optica l dens ity must be be low 0.1 at the excitation wavelength(350nm in this thesis). The selected excitation wavelength should be suffic iently far from the emission spectra. And it is advisable to use a single detector for all the spectra measurement to minimize the experimental errors introduced by detector positioning [START_REF] Yvon | A guide to recording fluorescence quantum yields[END_REF][START_REF] Laverdant | Experimental Determination of the Fluorescence Quantum Yield of Semiconductor Nanocrystals[END_REF].

-Absolute method In comparison with relative method, the absolute method directly obtains the quantum yie ld by detecting all sample fluorescence through the use of an integrating sphere, as illustrated in Figure 3.7. The inner surface of the integrating sphere is highly reflective to make sure no light leak out. The advantage of this method is applicable for both solution and film sample test. Note that due to the energy transfer between partic les in film sample, the quantum yield of film samples is usually lower than that of solution sample. Since, in our work, some large sized na nocrysta ls are not able to suspend stably in the solvent, absolute method will be a better way to determine the quantum yield. Absolute quantum yie ld of film samples was measured at the excitation wave lengths (410 nm and 450 nm) in this thesis, and the thin film samples were prepared by spin coating the particles on the glass substrate. The absorption and emission of glass substrate and sample are measured, respectively, and then quantum yield can be given using the ratio between emission photons and absorbed photons of sample as definition[119].

Transmission electron microscopy

Studying the morphology of the colloida l nanocrysta ls inc luding the shape and s ize distribution provides important information to verify the rationa lity and effic ienc y of synthetic protocol. Since the optical properties of the nanocrysta ls are also strongly related to the ir structura l properties, the complete understanding of the structura l properties of the samples is cruc ia l to the deve lopment and improvement of nove l heterostructura l nanocrysta ls w ith the we ll-defined s ize, sha pe and compos ition, leading to the achie vement of superior che mical and optica l properties. In this thesis , we have mainly used transmission electron microscopy (TEM) to investigate the structural properties of our synthesized nanocrystals. TEM is a useful technique to intuitively investigate morphology and crystalline structure of nanocrysta ls (Figure 3.8). Electrons beam in microscopy is created by an electron gun and accelerated under 200 kV high voltage, then it is transmitted through an ultrathin mesh copper grid, and the image is formed from the interaction of the electrons with the sample. The sample preparation process is very importa nt to the qua lity of the images, we usually first dilute our samples, the samples are necessary to be cleaned for several times to remove the excess amounts of organic ligands, which otherwise causes accumulation of charge at the surface under the electrons beam. Then, the diluted solution is dropped on the copper mesh grids and wa it for the evaporation of the nonpolar solvent. Fina lly, the samples are kept under vacuum overnight for the complete drying before the measurements.

Spherical quantum wells synthesis and characterizations

SQWs(CdS/CdSe/CdS) cons ist of thin CdSe emiss ive la yer sandw ished in two CdS layers. The strained epitaxial growth diminishes the misfit defects induced by the structura l strain, the high quantum yie ld of SQWs has been observed recently [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF] and its quas i type II bandgap alignment makes it poss ible to decrease the Auger effect by engineering the electron-hole wavefunction overlap. Therefore , SQWs are considered as a good candidate for investigation of the scintillation properties of nanocrystals. Additionally, owing to the release of the structural strain as the shell grows in SQWs, we expect to synthesize a series of SQWs on a large size scale without the loss of the chemical and physical qualities to achieve the study on size dependence of scintillating phenomenon in nanocrysta ls. The synthes is and characterization are introduced in the following sub-section. Unless otherwise specified, all chemical materials are used directly without further purification or treatment.

Preparation of precursors

Cd(myr)2. is prepared according to a procedure published previously [START_REF] Yang | Synthesis of CdSe and CdTe Nanocrystals without Precursor Injection Angew[END_REF]. 9.3g sodium myristate was dissolved in 750mL methanol and stirring until completely diss olved (~1h). And 3.7g cadmium nitrate diss olved in 120mL methanol was mixed into the above solution and stirred s low ly to obta in white precipitate(1h). Then the precipitate was filtered by a Buchner funne l, rinsed severa l times w ith metha nol a nd dried under vacuum overnight. 0.25M S-ODE. 20mL ODE was degassed under vacuum for 30 min at 70 °C to evacuate oxygen and water(P<0.3mbar), and the reaction was performed under argon gas condition and 160mg sulfur powder was added into flas k. Then the temperature was heated up to 140°C until all the sulfur powder was dissolve d, and the flask was kept at 80°C to prevent sulfur powder from being precipitated out of solution. 0.5M Cd(OA)2 in OA. 6.42g CdO was dissolved in 100mL ole ic acid at 180°C, when the mixture was slightly yellow and transparent, the temperature was decreased to 120°C to avoid the Cd(OA) 2 degrade at high temperature. Then the flask was degassed until P<0.3mbar. 0.5M Cd(OA)2 in OA/ODE. 6.42g CdO was dissolved in 50mL ole ic acid and 50mL ODE at 160°C for 30mins under argon, and the solution was degassed at 100°C for 90mins. Comparing to the 0.5M Cd(OA) 2 in OA, the existence of ODE can prevent the nanoparticles from degrading by oleic acid.

2M TOPSe. 1.58g Se powder was added in 10mL TOP at 100°C for 4h until the ir complete dissolution. The reaction was performed in glove box to exclude the influence of oxygen. Then TOPSe was stocked under inert gas atmosphere.

Synthesis protocol

Spherical quantum wells synthesis is a three steps process [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF]. CdS seeds are formed through homogeneous nuc leation, and followed by a thin CdSe emiss ive layer synthesis. The subsequent s hell growth is performed to cover different thic kness she ll on CdS/CdSe-based nanocrystals. For this subject, we are interested in samples coated by very thick s hell, thus continuous injection of s hell precursors was chosen to fulfill the formation of emissive layer and shell. The protocol is shown in Figure 3.9. The specific experimental steps and parameters are as follows:

Synthesis of CdS see ds. 1.5mL 0.5M Cd(OA) 2 in OA/ODE and 27mL ODE were quickly sw itched under vacuum and argon for three times. Under flow ing argon, the flask was heated up to 270°C, and 1.5mL S-ODE(0.25M) was injected at this temperature. Afterwards, temperature was decreased to 250°C immediate ly to undergo the growth process for 15mins. After cooling down the mixture to room temperature , the samples were purified by precipitat ion w ith etha nol, centrifugated at high speed and redis persed in hexane. The CdS seeds are around 3nm in diameter.

Synthesis of CdSe laye r. 2.7mL CdS seeds at 193μM and 5mL ODE were degassed under vacuum at 65°C for 1h. The prepared 0.1M precursors mixtures (0.6mL TOPSe, 2.4mL Cd(OA) 2 in OA/ODE and 9mL ODE) was injected at 6mL/h at 300°C. 11.2mL mixture was consumed for the des ired s ize emiss ion layer. The solution was coole d down to room temperature, and precipitated with ethanol, toluene and OA. The final sample was suspended in hexane.

Reaction equation of CdSe layer synthesis:

(3-7) Synthesis of CdS she ll. 3mL CdS/CdSe(52μM), 85mg Cd(myr) 2 and 10mL ODE are introduced in a three-neck flask. Cd(myr) 2 was added to prevent the degradation of CdS/CdSe at high temperature. And the additiona l Cd 2+ can invert the reaction as depicted in equation [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF][START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Diroll | Violet-to-Blue Gain and Lasing from Colloida l CdS Nanoplatelets: Low-Threshold Stimulated Emission Despite Low Photoluminescence Quantum Yield[END_REF][START_REF] Röntgen | On a New Kind of Rays[END_REF][START_REF] Dujardin | Needs, Trends, and Advances in Inorganic Scintillators[END_REF][START_REF] Létant | Semiconductor Quantum Dot Scintillation under γ-Ray Irradiation[END_REF].

Reaction equation of CdS/CdSe degradation:

(3-8) After 30mins of degassing under vacuum at 65°C, the solution was heated up to 305°C and 1mL OA m was injected to form monodis persed nanopartic les (It is better to inject after 2 drops precurs ors to avoid the CdSe layer da maged by OAm). Cd and S precursors were prepared as 0.2M solution of 1.58mL DDT diluted in 31.42mL ODE, and of 13.2mL Cd(OA) 2 in OA/ODE dissolved in 15.8mL ODE and 4mL OAm. The precursors were injected continuous ly into the flask by two syringes at 0.8nm/h under argon. Taking a liquots from flask at des ired s ize, 15nm, 20nm, 25nm, 30nm, 35nm, 40nm in diameter, the specific parameters are shown in Table 3.1. The aliquots were reheated up directly to 305°C, 1.5mL Cd(OA) 2 in ODE was added to passivate the surface of SQWs. After 30mins annealing, the reaction was cooled down to room temperature. The nanoparticles were prec ipitated and re dispersed w ith ethanol and hexane. Then the samples were suspended in hexane. 35nm and 40nm samples are unable to suspend in hexane due to the ir large size. For characterization, toluene has a good performa nce to disperse these large sized samples at low concentration.

Reaction equation of CdS shell:

(3-9)

Characterization

For II-VI semiconductor nanocrysta ls , most research has focused on the size of CdSe quantum dots rather than CdS [START_REF] Yu | Experimental Determination of the Extinction Coefficient of CdTe, CdSe, and CdS Nanocrystals[END_REF]120,121]. The published average diameter of nanocrysta ls is usua lly determined by transmiss ion electron microscopy image. Since the low accuracy of this method in the case of small size nanocrystal measurement, the size of CdS seeds in our work is calculated according to reference [START_REF] Maes | Size and Concentration Determination of Colloida l Nanocrystals by Small-Angle X-ray Scattering[END_REF] , in which a relatively accurate size is determined by mean of small-angle X-ray scattering(SAXS). The diameter is given from the following equation:

D = 1.12×10 -6 λ 3 + -1.13×10 -3 λ 2 + 3.81×10 -1 λ+(-40.5) (3-10)
where is the wavelength of the first-excitonic absorption peak. And the extinction coefficient can be calculated by the below equation [START_REF] Li | Extinction coefficient per CdE (E = Se or S) unit for zinc-blende CdE nanocrystals[END_REF]:

ε NC = 6.53× 10 -21 λ 10.11 [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF][START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Diroll | Violet-to-Blue Gain and Lasing from Colloida l CdS Nanoplatelets: Low-Threshold Stimulated Emission Despite Low Photoluminescence Quantum Yield[END_REF][START_REF] Röntgen | On a New Kind of Rays[END_REF][START_REF] Dujardin | Needs, Trends, and Advances in Inorganic Scintillators[END_REF][START_REF] Létant | Semiconductor Quantum Dot Scintillation under γ-Ray Irradiation[END_REF][START_REF] Hossu | X-ray luminescence of CdTe quantum dots in LaF 3 :Ce/CdTe nanocomposites[END_REF][START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF][START_REF] Chen | All-inorganic perovskite nanocrystal scintilla tors[END_REF] which is using the half area of the first excitonic absorption peak as the correction of size distribution. Based on the extinction coefficient, the molar concentration of CdS seeds is readily obtained by the Lambert-Beer law:

C = A ε NC l (3-12)
Here, A is the absorbance at the first excitonic absorption peak of CdS seeds in a given solution (hexane), l is the optical length (1 cm).

Absorption and photoluminescence spectra

The absorption s pectrum is performed w ith AvaSpec-2048L versatile fiber-optic spectrometer, which is used for measurements in the UV-VIS-NIR spectral range (200 to 1100 nm). A cuvette w ith pure hexa ne was measured as a reference, and then add ing some microliters of the samples into cuvette to acquire the spectrum of CdS seeds. As mentioned above, the s ize of samples can be determined by the pos ition of first exc itonic peak on absorption spectrum. We thus taken out aliquots from the flask during the reaction to monitor the growth of the nanoparticles. The final band edge absorption is centered around 392nm (Figure 3.10), according to the calculation, the size of these CdS seeds is around 1.3nm in radius. The growth of CdSe emiss ive layer is manifested on the spectrum, as shown in Figure 3.11(a), a liquots taken at different introduced precursors shows a red shift as the size increases, which is related to quantum confinement effect. We got over the reaction when the band edge absorption is at around 610nm. The size of CdS/CdSe is determined from TEM images, 6.1 0.5 nm in diameter in average (standard deviation of 8% compared to the average value), which corresponds to ~1.8nm thickness of CdSe emiss ive layer in Figure 3.11(b). This va lue is cons istent w ith the size we calculated before the synthes is. It means this layer grows homogeneous ly, all the precursors were consumed on the CdS seeds, rather than forming the secondary population of pure CdSe nuclei. The absorption and photoluminescence spectra of different size spherical quantum wells are presented in Figure 3.12. Photoluminescence spectrum is acquired on a FS5 spectrometer from Edinburgh Instruments. Samples sus pended in he xane or toluene (35nm and 40nm diameter SQWs samples cannot be dispersed in hexane) are excited at 405nm. In order to interpret the onset of the CdS she ll growth, we plot the CdS/CdSe spectrum as a reference. First, absorption at wavelength inferior to 510nm is dominated by the CdS shell, thus the increased absorbance indicates the growth of the shell. Meanwhile, the volume ratio of CdSe layer to CdS she ll reduce continuous ly, ma king that the absorbance of the CdS shell is prevalent and the peak features of CdSe are nearly disappeared and negligible. Secondly, an obvious red shift after the formation of she ll is attributed to the de loca lization of e lectrons in the entire nanopartic le. For a quas i-type II energy band alignment, the degree of confinement reduces as the size increases, leading to a continuous low energy shift. However, we observed the identical emission position for variety SQWs synthes ized, approximately at 650nm. The reason of this behavior is that we study ultrathick she ll nanocrysta ls , where the change in energy gap is a lready very small [START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF]. The FWHM of all the SQWs samples is around 30 nm, as we know, this va lue is strongly influenced by the polydispersity of the sample. Typically, CdSe quantum dots with a low polydispers ity have a FWHM of around 20 to 30 nm. Cons idering the difficult-to-control size distribution of our samples due to their large size, 30nm width basica lly corresponds to monodis persed samples. Compared to 55nm of FWHM in the gradient potential QDs with comparable dimension in the previous report [START_REF] Nasilowski | Gradient CdSe/CdS Quantum Dots with Room Temperature Biexciton Unity Quantum Yield[END_REF] , the interpretation of the narrow linewidth is the relatively smaller size dispersion, as well as the small structural strain dispersion in sandwich-like structure [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF]. All of the samples with same concentration under white light and UV light are shown in the Figure 3.13. Here, the quantum yie ld is measured with relative and absolute methods. The relative method is only applied for the measurement of smaller SQWs (15nm and 20nm). The quantum yie ld in solution are around 80% (reference sample is Rhodamine 101). And the quantum yie ld in film of larger SQWs (30nm and 40nm) are about 35% measured by abs olute method. Ow ing to the energy transfer in film sample , the quantum yie ld should be higher in the case of solution. This high quantum yie ld in such giant partic les suggests the influence of the misfit defects as the shell grows is negligible in SQWs compared to the conventional core/shell structure.

Photoluminescence excitation spectroscopy

Photoluminescence excitation spectrum is measured with low concentration samples (absorbance<0.1 at the whole measurement range) to exclude the impact of selfabsorption. As we introduced above, PLE spectra allow us to identify the secondary nucleation through the comparison with absorbance. We acquired the PLE spectra on FS5 spectrometer from Edinburgh Instruments. As shown in Figure 3.14, we compared PLE spectra and ABS of CdS/CdSe and spherical quantum wells. As shown in Figure 3.14(a), in comparison w ith ABS, the intens ity of PLE spectrum of CdS/CdSe behaves s lightly lower at short wavelength. After adding the carboxylate ligands, we observed an enhancement of the PLE intensity. Indeed, a part of the ligand is stripped during purification process, leading to the formation of the traps involved w ith dangling bonds. The capture of carriers by the traps is respons ible for the reduction of the PLE curve. And the additional ligands can decease the traps by bonding with the dangling bonds, resulting in the increase in the emission intensity. Therefore, the discrepancy between PLE and ABS is not caused by the secondary nucleation, we are allowed to conclude the samples are homogeneous.

For spherical quantum wells in Figure 3.14(b), we selected three emission wavelengths near the bandga p, the monodis pers ity of sample can be manifested with the identica l PLE. Besides, the intens ity of the PLE spectrum below 510nm should ma inly indicate the presence of CdS shell. However, when compared with ABS, a discrepancy between two spectra can be seen, implying a sma ll part of secondary nuc le i CdS core existence in this sample. The impact of this small quantity of nuclei is acceptable in our work. And size-selected purification can be applied to separate two populations with different size.

Structural characterization

Figure 3.15 is the transmission electron microscopy images with different sized SQWs, the measurement was performed at magnification 30000X. The sample with a relative ly low concentration was dropped on the copper grid. And then this mesh grid was degassed under vacuum overnight to remove liga nds and solvent. TEM images allow us to visually compare changes in size and shape of samples. We clearly observed the increase in nanopartic le s ize , implying the successful she ll precursors depos ition on the core. The synthesized particles exhibit a faceted morphology and somewhat irregular shapes, as frequently observed for large core/shell QDs [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF]. Since the shape of our samples is strictly not spherical, we selected maximum Feret diameter to represent the size of these samples. The size is estimated by software ImageJ, but the overlap between nanoparticles in large sample image compe ls us to measure them manually. The diameters of the samples are characterized from 15 nm up to approximately 40nm, exceeding the size of almost all published QDs until now. The size distribution standard deviations are around 12%, 10%, 15%, 10%, 10% and 12%, respective ly (Figure 3.15, inset). These size dispersions are well enough for such large nanoparticles. We have synthes ized a series of SQWs(CdS/CdSe/CdS) w ith different she ll thickness, the total diameter is from 15nm up to around 40nm. The successful coating of the shell is demonstrated by a nalyzing the optica l spectra and TEM morpho logy, a ll the samples emit at approximate ly 650nm w ith a narrower FWHM(~30nm). The narrow linewidths and about 10% deviation from TEM images are the evidence for the good size dispers ion. Additiona lly, we have prove d the SQWs synthesized exhibit high quant um yield ~80%, even for the giant 40nm SQWs the quantum yield in film is around 35%, which is the result of the reduction of misfit defects due to the re lease of the structura l strain in SQWs.

Colloidal Synthesis and Characterization of Hererostuctural Nanoplatelets

Nanoplate lets ha ve been applie d in ma ny applications s ince it was first synthesized in 2008 [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF]. In this thes is, we will discuss another potentia l application as a kind of scintillating material. Due to the giant oscillator strength in nanoplatelet, its decay time is much faster than in traditional QDs, which is a significantly attracting property to the fast-timing radiation detection. Therefore, here, our a im is to synthes ize different heterostructural NPLs: core/crown(C/C), core/shell(C/S) and core/crown/shell (C/C/S) to explore the underlying phys ics under X-rays excitation, and also to compare the difference on their scintilla t ion behaviors.

Chemical materials

-1-octadecene, 90%,CAS : 112-88-9. Sigma Aldrich -1-dodecanethiol, 98%, CAS : 112-55-0, Sigma Aldrich -cadmium acetate dihydrate, >=98%, CAS : 5743-04-4, Sigma Aldrich -cadmium nitrate tetrahydrate, >99%, CAS : 10022-68-1, Sigma-Aldrich -ethanol,>=99%, CAS : 64-17-5, Carlo Erba -hexane, CAS : 110-54-3, Carlo Erba -oleic acid, 90%, CAS : 112-80-1, Alfa-aesar -oleylamine, 80-90%, CAS : 112-90-3, Acros -selenium, 99.99%, CAS : 7782-49-2, Strem -sulfur, 99.99%, CAS : 7704-34-9, Strem -sodium myristate, >98%, CAS : 822-12-8, TCI -toluene, 99.9%, CAS : 108-88-3, VWR -trioctylphosphine , 97%, CAS : 4731-53-7, Sigma-Aldrich -zinc acetate dihydrate, >=99%, CAS : 5970-45-6, Sigma Aldrich Unless otherwise specified, all chemical materials are used directly without further purification or treatment.

CdSe nanoplatelets

First, we started with the synthesis of CdSe core NPLs: 4ML CdSe nanoplatelets having four complete monolayers of CdSe w ith an additiona l layer of Cd atoms were synthesized with a slightly modified recipe from the literature [START_REF] Ithurria | Quasi 2D Colloidal CdSe P latelets with Thicknesses Controlled at the Atomic Level[END_REF]. In a three-neck 250 mL flask, 1.02g of cadmium myristate and 70 mL of ODE were loaded. The solution was degassed and mixed under vacuum at 90°C for 1 hour, the vacuum pressure being lower than 0.3mba r, to remove oxygen and undes ired solvents. Then, 72 mg Se was dissolved in 2mL ODE, a nd it was rapidly injected into the flask under an inert atmosphere (argon). Afterwards, the temperature of the solution was set to 240°C. When the solution color becomes bright yellowish (generally around 205°C), 300mg cadmium acetate was added. The time of acetate addition is very important for the reaction yield of 4ML NPLs. After reacting for 10 min at 240°C, the solution was cooled to 150°C and 6mL OA was injected. Adding OA at high temperature is more effective, while injection of OA at low temperature will cause the aggregation of NPLs. Then, the temperature was continued decreasing to RT. When the solution temperature approached to the boiling point of hexane 69°C, 90 mL of he xane was injected to avoid the stacking of NPLs. The preliminary solution usua lly conta ins the surplus precursor and the byproduct as 3ML NPLs and QDs. Due to the difference in the size of different types of particles, we separate 4ML NPLs through size-selective purification step. By adjusting the ratio of solvent and anti-solvent, 3ML NPLs can be centrifuged out first. Next, further dissolution and re-precipitation can separate out ultra-small QDs. Therefore, we can get rid of the undesired byproduct and obta in the pure 4ML NPLs. Fina lly, the nanoplate lets were dispersed in hexane. Absorption and photoluminescence emission spectra of these CdSe core NPLs are given in Figure 3. 16(a,b). The sharp excitonic features in absorption spectrum are observed around 480nm and 511nm, corresponding to e lectron/light hole (LH) a nd electron/heavy hole (HH) transitions, respectively, and the feature around 420nm is assigned to split off (SO) transition [START_REF] Ithurria | Colloida l nanoplatelets with two-dimensional electronic structure[END_REF]. Furthermore, the very sharp emission profile of NPLs at 512nm is observed in the PL spectrum, the linew idth is narrower than 10 nm, indicating the high monodispersity. TEM ima ge of CdSe NPLs is show n in Figure 3.17. Although the morphology is irregular, most of the particles are rectangular-like, with the smallest lateral side length about 10nm. And the thickness of 4ML is around 1. 2nm. So both of side lengths surpass the exciton Bohr radius of 2D CdSe , which is around 3.5nm [START_REF] Tessier | Spectroscopy of Single CdSe Nanoplatelets[END_REF]124], only vertical dimens ion is confine d and no lateral confine ment could be observed. These CdSe NPLs will be used as a core for the subsequent synthes is of core/crown, core/she ll and core/crown/shell architectures.

Core/crown CdSe/CdS

In the past decade, the core/crown NPLs structure has received tremendous attention due to its excellent phys ical and che mical properties [START_REF] Tessier | Efficient Exciton Concentrators Built from Colloida l Core/Crown CdSe/CdS Semiconductor Nanoplatelets[END_REF]125,[START_REF] Leemans | Near-Edge Ligand Stripping and Robust Radiative Exciton Recombination in CdSe/CdS Core/Crown Nanoplatelets[END_REF]. Here, we tried several different approaches to perform the synthesis , in order to obtain high qua ntum effic iency and uniform shape , inc luding shell precurs ors mixture solution injection and different kind of anion precursor injection.

Approach 1: Anisotropic Growth Solution injection.

Anisotropic growth solution was prepared follow ing the previous ly publis hed procedure w ith s light modifications(scale up the qua ntity of precursors) [START_REF] Tessier | Efficient Exciton Concentrators Built from Colloida l Core/Crown CdSe/CdS Semiconductor Nanoplatelets[END_REF]. The preparation is divided into two steps. First, for Cd precursor, 960 mg of cadmium acetate dihydrate, 680μL of OA, and 4 mL of ODE were loaded in a three-neck flask. The solution was heated to 100°C under vacuum for 20 min until the formation of whitish color homogeneous ge l. For the preparation of S precursor (0.1M S-ODE), 150mL of ODE in 250 mL three-neck flask was degassed to remove a ir and water at 70°C for 30 min. Then under argon 480 mg s ulfur was added. The temperature was set to 140°C until a slight ye llow s olution formed and then cooled down to RT. Then Cd precurs or was mixed w ith 6 mL S precursor, heating and sonication steps were followed alternatingly till two precursors completely mixed, and then the anisotropic growth solution was used for the CdS crown coating.

For crown synthes is, 3mL 4ML NPLs (20 μL core nanoplatelets dissolved in 3mL hexane has an absorbance of 0.15 at 350nm) and 5mL ODE were loaded into three-neck flask. The hexane was removed in vacuum for 1 h at 65°C. Subseque ntly, the temperature was heated to 210°C under argon and then different quantities of anis otropic growth solution was injected at 8 mL/h. After the injection, the solution was cooled down to room temperature and 1mL OA was added at 80°C to avoid the aggregation. The solution was purified by ethanol and hexane only once. Fina lly, the C/C NPLs was dispersed in hexane.

Approach 2: Single precursor injection--S-ODE

First, a 0.05M S-ODE precursor was prepared by mixing 8mg of s ulfur pow der into 5mL of ODE. Afterwards, it was ultrasonicated for at least 30 min until a s lightly turbid solution w ithout vis ible precipitate. For crown growth, 4m L of ODE, 190uL of OA a nd 160 mg of Cd(Ac) 2 were combined into a 50 mL three-neck flask, then degassed for 1h at 65°C. 5mL NPLs was injected into the flask under vacuum. Then the solution was heated up to 220°C under argon. When the temperature reached 210°C, the S precursor was injected at 20 mL/h. After the injection of desired amount precursor, the solution was cooled dow n to room temperature, and 0.6mL OA was added at 90°C to avoid the aggregation. The solution was purified by ethanol and hexane only once, and then dispersed in hexane.

Approach 3: Single precursor injection--TOPS

Alternative ly, crow n synthes is based on the TOPS was a slightly modified procedure from the literature [125]. At first, we prepared 1 M TOPS solution by dis persing 160.3 mg of sulfur into 5mL of TOP. The solution was stirred overnight at room temperature in an air-free glovebox to ensure the S powder was complete ly dissolved. To avoid the oxidization, the solution was stored in the glovebox. Afterwards, 0.5mL of 1M TOPS solution was diluted in 9. 5 mL ODE to form 0. 05M TOPS, preparing for the crown growth procedure.

To overcoat a CdS crown, 8 mL of ODE, 180 mL of OA, 96 mg of Cd(Ac) 2 and 4mL of 4ML CdSe NPLs were combined into a 50 mL three -neck flask. As usua l, the mixture was degassed in vacuum for 1 h at 60°C. Subsequently, the solution was heated up to 240°C under an argon atmosphere. When the temperature reached 215°C, 2 mL of 0.05M TOPS precursor solution was injected at the rate of 12 mL/h to form a thin crown. After completion of the injection, unlike the previous two approaches, the solution was directly cooled dow n to room temperature w ithout additiona l OA. The solution was washed by addition of a 2: 1 ethanol : hexane mixture. The resulting C/C NPLs was finally redispersed in hexane.

The absorption and photoluminescence spectra of C/C NPLs synthes ize d by different methods were presented in Figure 3.18. They exhibit consistent electronic features. In comparison with core 4ML NPLs, with the formation of CdS crown layer, a new absorption peak appears at around ~400nm. By adjusting the amount of precursors, the CdS crown absorption peak becomes more pronounced, leading to a significant increase in the absorption cross section. However, the PL spectra show no significant change compared with core NPLs, this is ascribed to the lateral growth, maintaining the confine ment on vertica l dime ns ion. The emiss ion peak of all synthesized C/C NPLs is at approximate ly 515nm, only a slight redshift (around 3nm) in the spectra can be observed, which is caused by the change of die lectric constants due to the formation of the CdS crown[127]. From the TEM ima ges, the core/crown NPLs synthesized w ith S-ODE involved approaches usua lly show the irre gular shape, as shown in Figure 3.19(a) and (b). On the contrary, the particle fabricated with TOPS gives a regularly shaped C/C NPLs with straight edges (Figure 3.19(c)) and this extreme uniform shape has been reported to cause a further increase of the quantum yield. It is also certain that the regular shape is helpful for subsequent overcoating and application. Replac ing the common S-ODE with TOPS was demonstrated to affect the reactivity of sulfur, and the low reactivity was responsible for the formation of uniform shape [125]. Additionally, the PLE of C/C NPLs measured at three different selected wavelengths is illustrated in Figure 3. 19(d), the emission at these three wavelengths originates from the same luminescent species since exactly similar spectral profiles are obtained, suggesting a great monodispersity in this core/crown NPLs. The analogous PLE has been observed for all the C/C NPLs in this thesis .

Hot-injection alloyed core/shell CdSe/CdZnS

Motivated by the core/shell structure of QDs, in order to improve the performance of NPLs, different approaches and materia ls have been constantly investigated to fabricate heterostuctura l core/she ll NPLs. To date, shell coating is usua lly achieved by two approaches. First, by a lternate ly depos iting cation and anion layers at low temperature, colloidal atomic layer depos ition (c-ALD) techniques are able to form highly uniform growth she ll w ith atomica lly precise thickness control [START_REF] Ithurria | Colloida l Atomic Layer Deposition (c -ALD) using Self-Limiting Reactions at Nanocrystal Surface Coupled to Phase Transfer between Polar and Nonpolar Media[END_REF]. However, c-ALD approach usually suffers from low quantum yield and poor stability issues. On the other hand, by adding the precursor of shell at high temperature, one can obtain core/shell NPLs with uniform and high quantum efficie ncy. But the lower therma l sta bility of NPLs, compared to that of QDs makes the high-temperature shell growth method difficult to perform in the NPLs system. To prevent etching and dissolution of the NPL cores prior to she ll growth, Rossinelli et al. investigated different reaction solvents to deve lop a hot-in je ct ion shell growth method to coat a thick CdS she ll on CdSe NPLs [START_REF] Rossinelli | High-temperature growth of thick-shell CdSe/CdS core/shell nanoplatelets[END_REF] , which was demonstrated with predictable better chemical and optical quality. Hence, our synthesis is according to the recently reported protocol to form an alloyed CdZnS by hot injection [START_REF] Altintas | Giant Alloyed Hot Injection Shells Enable Ultralow Optic al Gain Threshold in Colloida l Quantum Wells[END_REF]. Specifically, 14mg of Cd(Ac) 2 , 34mg of Z n(Ac) 2 , 0.5 mL of OA, 5 mL of ODE and 5mL of CdSe core NPLs was loaded in a 50 mL threeneck flask together. The mixture was degassed at 80 °C for 1h to remove oxygen, moisture, and hexane. Meanwhile, the precursors were reacted to form Cd(OA) 2 and Zn(OA) 2 . Then, the solution was heated up to 300°C under argon atmosphere. Subsequently, 0.5mL of ole ylamine was injected quickly at 90°C. 350μL of 1-Octanethiol dissolved in 20mL of ODE was injected at 165 °C at 10 mL/h. The addition of precursors at relatively low temperature can limit the destruction/etching of the core NPLs. Then the injection rate was decreased to 4 mL/h when the temperature reached 240°C. After the complete injection, the solution was kept for 50 min at 300°C for annea ling and full consumption of unreacted precursors. Then the solution was rapidly cooled down to room temperature, and 5 mL hexane was added at 60°C. Fina lly, the core/she ll NPLs was precipitated by ethanol, and redispersed in hexane. We successfully synthes ized three different thic knesses core/shell NPLs, emitting at 615nm, 631nm and 671nm respectively. The absorption and photoluminescence spectra are given in Figure 3.20. In comparison with core NPLs, the absorption spectra also possess the clear HH and LH transitions peak, but the difference is that the excitonic features have been experienced a strong red shift when the shell is coated. This red-shift continuously increases with the increase in CdZnS shell thickness which can be explained by the decrease in quantum confinement. Besides, since the Cd reactivity is higher than that of Zn, the inner shell is mainly composed of CdS, and the outer surface is Zn-rich. The conduction band a lignment difference of CdSe and CdS is ne gligible. Electron delocalization into the shell could be another reason for the excitonic features redshifting. In addition, as the shell increases, the absorbance in the blue spectral region is enhanced, resulting in the larger absorption cross-section. And the FWHM of PL spectra are around 25nm for a ll the samples. The broadening, compared to 4 ML NPLs, is attributed to stronger coupling to phonons in the shell [START_REF] Cui | Evolution of the Single -Nanocrystal P hotoluminescence Linewidth with Size and Shell: Implicatio ns for Exciton-Phonon Coupling and the Optimization of Spectral Linewidths[END_REF]. The quantum yie ld in thin films is measured as approximately 50% for all the samples by absolute method. In order to verify the uniformity of the she ll growth, we measured the core/she ll CdSe/CdZnS NPLs PLE spectra at different emiss ion wavele ngths, as shown in Figure 3.21. The same excitonic features taken at fluorescence peak position and on both sides have been observed, which suggests a uniform coating of the she ll a nd the induced inhomogene ous broadening is negligible. Moreover, from the TEM images in Figure 3.22(a), the shell growth is manifested as the change of the vertical thickness of nanoparticles , the thickness of the sma ll core/she ll NPLs is determined as around 3.5nm (4ML core NPLs thic kness is ~1. 2nm) , which matches with the estimated thickness of ~1.2 nm. It corresponds to 4 monola yers shell. A nd the thickness for medium and large core/shell NPLs are estimated as 4.2nm and 9nm, corresponding to 5ML a nd 12ML she ll. Though the core shape is irregular, all the core/shell NPLs appear the rectangular shape, as presented in More complex architecture core/crown/shell NPLs has been recently employed to achieve amplified spontane ous emiss ion(ASE), w ith an ultra low threshold of ~ 800 nJ/cm 2 . This is attributed to the combination of the characteristics of C/C and C/S in this structure, that is, large absorption cross-section and suppressed Auger non-radiative recombination [START_REF] Taghipour | Sub-single exciton optical gain threshold in colloidal semiconductor quantum wells with gradient alloy shelling[END_REF]. However, so far, the synthesis approaches of C/C/S NPLs are all based on c-ALD [START_REF] Kelestemur | Platelet-in-Box Colloidal Quantum Wells: CdSe/CdS@CdS Core/Crown@Shell Heteronanoplatelets[END_REF][START_REF] Shendre | Ultrahigh-efficiency aqueous flat nanocrystals of CdSe/CdS@Cd 1-x Zn x S colloida l core/crown@alloyed-shell quantum wells[END_REF][START_REF] Taghipour | Sub-single exciton optical gain threshold in colloidal semiconductor quantum wells with gradient alloy shelling[END_REF] , resulting in low quantum yield and poor stability. In this thesis, we try to apply the hot-injection protocol to fabricate the core/crown/alloyed she ll CdSe/CdS/CdZnS NPLs.

The procedure is similar with C/S synthetic method. Core NPLs were replaced with 3mL C/C NPLs(10mL in tota l, s ynthesized w ith TOPS) as the seeds, and they were mixed with 12 mg of Cd(Ac) 2 , 27.5 mg of Zn(Ac) 2 , 0.5 mL of OA and 5 mL of ODE into 50 mL three-neck flask, follow ing by the degassing and reaction process. Then the solution was heated up to 300 °C under argon atmosphere. 0.5mL of oleyla mine was injected quickly at 90°C. 350μL of 1-Octanethiol was dissolved in 20mL of ODE, and the injection rate was set as the same as above 10mL/h first, and the n turned to 4mL/h until 7mL precursor was consumed. After the complete inject ion, the solution was kept for 20 min at 300 °C for annealing and then rapidly coole d dow n to room temperature. After washing and precipitating, the final C/C/S NPLs was dispersed in hexane.

In the Figure 3.23(a), we presented the absorption and photoluminescence spectra of C/C/S NPLs and an absorption spectra of C/S NPLs with the similar emiss ion wave length (dark dash). Except for the heavy hole and light hole transitions, the C/C/S exhibits a larger absorbance than that of comparable C/S NPLs, particularly in the blue spectral region. The marked increase at around 500 nm that is related to absorption from the coating CdS crown. However, the exc itonic features are broadened, the width of the emiss ion profile is around 40nm, larger than that of the regular C/S NPLs FWHM. To examine the uniformity of our sample, the PLE is shown in Figures 3.23(b): the spectra at different wavelengths of PL peak show the similar features, indicating that the origin species are not very distinctive. Therefore, the polydispersity of the particles is not the reason for the broadening of the spectrum From the TEM images in Figure3.24, the lateral shape is nearly monodispersed and rectangular with average size (16nm and 34nm). The thickness in the center is around 3nm, corresponding to 3ML she ll. However, the larger contrast has been observed at the corners, indicating the vertical size is thicker at the corner than the center. The resulting C/C/S NPLs thus form overall bone-shaped nanocrysta ls. Actua lly, this morphology has also been observe d in core/she ll. The presuma ble growth process was reported by Andrea Castelli et al. [START_REF] Castelli | Core/Shell CdSe/CdS Bone-Shaped Nanocrystals with a Thick and Anisotropic Shell as Optical Emitters[END_REF]. They explain it as anisotropic shell with double structures(zinc blend and wurtzite type). A cubic she ll at first grows in thic kness, and then {111} facets might be available at the corners, which provide suitable locations for the growth of a wurtzite shell. Here, I think the efficient change observed in C/C/S is due to the peripheral CdS crown promotes the wurtzite she ll growth. A nd this bone shape should be respons ible for the larger linewidth. The quantum yield in thin film was measured as ~25%, the reduction of the quantum yie ld compared to that of C/S NPLs is unknow n, one of the interpretation might be the defects generated on the corners. We have synthesized three different types of heterostructura l nanoplate lets. First, for core/crown CdSe/CdS NPLs, by us ing the TOPS instead of the S-ODE as a S precursor, we achieved the improvement of the morphology with the rectangular shape , which is beneficial of the subsequent coating and the application. The low reactivity of sulfur in TOPS has been demonstrated to be respons ible for the formation of uniform shape [125]. Secondly, we prepared the core/shell CdSe/CdZnS NPLs with three varying thicknesses by hot-injection approach. The quantum yie ld in film was kept around 50% for all the samples, which indicated the alloyed shell leads to the decrease in the defects induced by the lattice mismatch. Therefore, we will employ these samples to investigate the size dependence of sc intillation properties on 2D nanocrysta ls. Finally, for the first time, we synthesized the core/crown/she ll NPLs via the hot-injection methods, 3ML CdZnS was successfully coated on the CdSe/CdS. Though the broa denin g emiss ion linew idth a nd a moderate quantum yie ld in thin film were obta ined, due to larger absorption cross section, we believe that with the further deve lopment on the synthes is, this configuration w ill be promising.

Conclusion

In this section, based on the funda menta l colloid nanocrysta l synthetic theory, we first successfully synthesized a series of SQWs with different shell thickness. The total diameters of resulting SQWs are from 15nm up to approximately 40nm in diameter, exceeding the size of almost all published QDs until now. Thanks to the release of the structural strain in this sandwich-like co nf ig urat io n , t he misf it defects in g ia nt part ic le s is much less than in conventional QDs, which is demonstrated by the high quantum yield (up to 80%) for 15nm and 20nm SQWs, even for the larger partic les(40nm) the quantum yield is around 35% in thin film. The samples are also characterized with narrower photoluminescence spectra linew idth and an homogeneous morphology, suggesting the good quality of the synthesized samples. Besides, we have also fabricated different heterostructural NPLs, including core/crown, core/shell and core/crown/shell. Through using TOPS replaces the traditional S-precursor, we achieved the formation of the regular shape of core/crown NPLs. And via hot-injection methods, three different sized core/shell NPLs with high qua ntum yie ld have been prepared. For the first time, we synthes ized a core/crown/she ll NPLs through the hot-injection method, though the quantum yie ld is only 20% in film, this experiment proved the possibility to fabricate this structure by hotinjection methods. These samples will be applied in the subsequent research regarding scintilla t io n mechanism of nanocrystals under ionizing radiation.

Chapter 4

Zhu MENG 79 Chapter 4

Scintillation properties of spherical quantum wells under X-rays excitation

The breakthrough of the Auger recombination suppression has stimulated the development regarding the application of the multiexc itons phenomena in semiconductor nanocrysta ls , such as low-threshold NCs lasers, light emitting diodes, as well as solarenergy convers ion [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Klimov | Single-exciton optical gain in semiconductor nanocrystals[END_REF][134][START_REF] Lim | Influence of Shell Thickness on the Performance of Light-Emitting Devices Based on CdSe/Zn 1-X Cd X S Core/Shell Heterostructured Quantum Dots[END_REF]. In addition, another attractive application is for radiation detection scintillator materia ls. So far, there have been several related reports [START_REF] Létant | Semiconductor Quantum Dot Scintillation under γ-Ray Irradiation[END_REF][START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF][START_REF] Chen | All-inorganic perovskite nanocrystal scintilla tors[END_REF][START_REF] Dujardin | Luminescence and Scintilla t ion P ropertie s at the Nanoscale[END_REF], but a few were reported on the energy depos ition and re laxation process of high-energy photons in NCs.

Therefore, in this chapter, we first introduce several characterization techniques used here, including steady-state radioluminescence spectroscopy and time-resolved measurement. Next, we will compare the temporal and spectral features under intense optica l excitation and X-rays excitation of SQWs with different sizes to realize the identification and assignment of excitation species, involving single exciton and multiexcitons recombination. The scintillating behavior evolution with the size of SQWs is further given. By comparison of spectral ana lys is w ith numerica l s imulations , we are able to draw a picture of a fundamenta l understanding of the underlying phys ics in the scintilla t io n process of such nanocrystals.

Experimental setups 4.1.1 Radioluminescence setup

The light emitting induced by ionizing radiations is referred as scintillation sometimes also called radioluminescence(RL). The measurement setup under X-rays excitation in our laboratory is presented in Figure 4.1, which cons ists of an X-rays tube, a monochromator and a CCD. The technical difference from PL measurement setup is that the inc ident X-rays is generated in a protective cover to prevent the leak of the radiation. And then for the sample preparation, unlike that of the photoexcitation, glass and quartz substrate are usually not the good candidates for measuring scintillation luminescence under X-rays excitation, because the excitation can be strongly absorbed and their intrinsic scintillation signal, even with a poor yield introduce a significant effect on the sample spectrum because of the significant X-rays absorption as compared to the studied film (a few tens of nm). Furthermore , the amount of nanopartic les in the solution sample is re lative ly low, making an ineffic ient absorption and emiss ion intens ity under X-rays radiation. Hence, instead of dispersed NCs in cuvette, X-rays induced luminescence can only be performed us ing a solid film sample dropped on Si substrate. Additiona lly, our setup can also be employed to achieve the PL measurement, facilitating the spectra comparison with radioluminescence.

Time-resolved measurements

Time-resolved spectroscopy is a useful technique to study excited state dynamics of nanocrysta ls and the understanding of complex spectral shapes. It gives us the opportunity to observe, explore and understand the fast time multiexc itons events taking place in the NCs. Here, decay curve measurement and time resolved emission spectrum are presented.

Decay curve measurement

Decay curve measurement is using time-resolved spectroscopy to record the fluorescence decay curves of the samples. Since the PL signal is proportional to the average number of exciton formed in the excited NCs, it allows to directly measure the exciton population in time follow ing the different kinetic recombination processes. The setup cons ists of a pulsed laser emitting at 405nm, a monochromator, a photomultiplier tube (PMT), PicoHarp300 time-correlated s ingle photon counting (TCSPC) and several filters , as show n in Figure 4.2. In principa l, a pulsed laser excites the sample a nd triggers the counting, and then the emitted photons from the sample after a certain time are detected by a PMT at a selected wavelength. The electronic system records the time between the laser pulse generation a nd the sample's photon detection. At the end, the detected photons are counted as a function of the time. This process is repeated millions of times. These photon counts from many exc itation events compose a his togram to form decay curve over time. The smallest time bin of our TCSPC is 16ps, while the overall full width half max spans from 200ps to 400ps depending on the set power of the laser, w hich is high enough to ana lyze the multiexc itons w ith a lifetime of on the order in the ns range. Moreove r, the laser repetition rate must be slow enough to ensure that all the excitonic species have time to recombine before the next laser pulse. Typically, the repetition time is at least 3 times longer than the longest lifetime of the NCs. In addition, in order to avoid artefacts caused by the high emission intensity, we keep the input counts value lower than 5% of the laser repetition rate. By analyzing the kinetics of fluorescence decay curves, one can obtain information on the sample's photoluminescence lifetime, which is an important characteristic of NCs. Theoretically, this decay curve can be fitted by an exponential with a time constant = 1⁄ whic h is the lifetime:

( ) = (0) = (0) (4-1)
However, when cons idering the polydis persity of the sample or severa l recombination pathways, like the multiexcitons, these yields to different monoexponential decays, resulting in the multiexponential fitting of the measured decay curve as in the following equation:

+ + + + (4-2)
By adjusting neutral density filter, we are allowed to acquire the decay curve under variety laser powers. Based on the definition of the multiexc itons , we can identify the excitonic species through the analysis of the power dependence on each time constant. Therefore, it is a very commonly used technique to study the radiative and non-radiative recomb inat io n of mu lt ie xc it ons. Besides, in th is t hesis , th is setu p c an be mod if ie d to perform the decay curve under X-rays excitation. The specific adjustments can refer to the relevant part in Figure 4.3.

Time resolved emission spectrum

Due to the ultra-fast recombination characteristics of multiexcitons, steady-state spectroscopy is not the most effic ient approach to observe multiexc itons emiss ion features. Inspired by the time-resolved decay curve measurement, we take the decay curves at different wave lengths , and then integrate the counts of e mitted photons at different times over the whole measured spectral range through a homema de progra m to obta in fluorescence spectra under different time windows, which is called as time resolved emission spectrum. In this case, we are allowed to extract the extremely fast recombination components from the entire spectrum, enabling a direct comparison between multiexcitons and single exciton recombination spectra. Combined with the fluorescence decay curve, it can be beneficial for us to have a deeper understanding of the excitons dynamics in nanocrystals. The schematic representation of experimenta l setup is given in Figure 4. 3, w hich is exemplified as X-rays exc itation. It cons ists of a pulsed laser, an X-rays tube, a PMT, a monochromator and a multic hanne ls counting MCS6A. Sa mples were mounted in front of the w indow of the X-rays tube. First, a laser pulse triggers the counting to start, simultaneous ly, pulsed X-rays w ith same repet ition rate are generated by the laser at 405nm to excite the samples. Then by scanning the wavelength in monoc hromator, the emitted photons are collected by PMT after a certain time to form the decay curve of different wavelengths in MCS6A. Subsequently, the home made Labview program integrates decay curves for all wavelengths describing the emission spectrum. The time gating is done after a post treatment, w ith a variety of time windows, giving a time resolved emission spectrum. Similar to the lifetime measurement, it can also be used for photoexc itation measurement w ith s light modification (Figure 4.2). The power dependent spectra can be acquire d by adjusting the neutra l dens ity filter. However, the difference is that the TCSPC is replaced by a multichannel counting, to overcome the artificial issue caused by the deadtime of TCSPC in the case of a great amount of emitte d photons. The drawback is the minimum time bin restricted to 800ps, making it less resolving for the lifetime record.

Possible mechanism in nanocrystals and simulation

P ossible scintilla tion process in nanocrystals

It is well know n that multie xc itons can be produced by absorbing more than one photon in a very short time window from the same pulse upon intense optica l excitation, follow ing by the therma lization process and relax the carriers to the lowest energy state, as shown in Figure4. 4(a). Alternative ly, the multie xcitons can also be generated through a carrier multiplication process with the absorption of a single photon whose energy is at least twice than the energy gap. Twice the gap is the minimum required from energy conservation cons iderations, but in genera l it is admitted that it require s several times the gap [START_REF] Klimov | Multicarrier Interactions in Semiconductor Nanocrystals in Relation to the Phenomena of Auger Recombination and Carrier Multiplication[END_REF]. The hot electron releases its energy by inelastic scattering with another valence electron, promoting it into the conduction band. Ins pired by the carrier multiplication of moderate energy photon, the involved mechanis ms under X-rays irradiation and more generally under ionizing radiation can be speculated. The energetic electron interaction with matter is undoubtedly a complex process involving multiple time, energy and space scales, therefore the process differs significantly from the moderate photon exc itation. In brief, the X-rays photon interaction with materials leads to the photo-ionization of an electron from a core band generating a primary highly exc ited electron-hole pair composed of a core hole and a hot electron (several keV). The primary electrons essentia lly escape from the materia l and generate secondary e lectrons by e lastic collision. Subsequently, these unstable hot carriers will undergo the carrier multiplication process. Once secondary e lectron-hole pairs cannot be produced due to the lack of energy as compared to the gap, the thermalization process becomes dominant and phonons are produced, ultimately forming exc iton and multiexc itons and emitting light when the radiative transitions are favorable. The plausible mechanisms are sketched in Figure 4.4(b). The hole relaxes ma inly through Auger cascade generating as well secondary electron-ho le pa irs. In a le ss extend , X-rays florescence occurs. Unlike bulk scintillating materials, instead of transporting the charge carriers up to their radiative recombination or their energy transfer to a luminescence center [START_REF] Nikl | Scintilla t ion detectors for x-rays[END_REF][START_REF] Rodnyi | Physical processes in inorganic scintillators[END_REF] , the case on nanocrysta ls experience significant escape of charges, depending on the particle size, as described in Bulin et al. [START_REF] Bulin | Modelling energy deposition in nanoscintillators to predict the effi-ciency of the X-rayinduced photodynamic effect 9[END_REF]. Motivated by this hypothetical analysis, we will explore the scintillation characteristics in NCs by comparing the optica l response under X-rays and intense light excitation. Excitation of NCs under X-rays is due to fast electrons created from primary photoelectric event in some SQWs. A fast electron with kinetic energy from 10 to 30 keV produces secondary e lectronic exc itations (SE) in a NC. The first part of the s imulation estimates the number of produced secondary excitations as a function of the particle diameter. These secondary excitations (SE) have different tota l energies, and the distribution of SE energies in bulk materia ls is proportiona l to the energy loss function (-1/ (ℏ )) where ε is the complex dielectric permittivity of the material, and it depends on the energy lost ℏ . It can be obtained from the Livermore database. For isolated NCs the distribution of SE over energies deviates from the bulk energy loss function due to appearance of surface plasmon resonances, but, for packed NCs, we can use the bulk approximation. Because of the complexity, we have considered a two-step procedure. First, the distribution of the number of scattering events per particle for the primary electrons forming 1, 2, 3... energetic secondary e lectron-hole pairs are estimated. Then, for a given number of scattering, i.e. initial number of SE, we estimate the time dependence of the population of the formed neutral excitons (X, XX, XXX...) as well as charges species (X + , XX + ...). This second part takes into account the electron-electron interaction, the thermalization i.e the electron phonon scattering, the Auger process, a work function of W=5eV as well as the radiative recombination cons idering the following cascades (case of XXX) XXX -> XX + ℎ -> X+ ℎ -> ℎ . The simulation is based on rate equations w ith changing of the number and charge of excitations. It supposes a strong interaction of excitations in NC; the coeffic ients of rate equations are proportional to the densities of states. As for the distribution of the number of electron-electron scattering for the primary electrons, we can approximate the mean energy of these excitations to be about the plasmon energy corresponding to the maximum of the energy loss function presented in For clarity, we present the time evolution of populations of neutral spec ies for a 25nm partic les in Figure 4.7. In order to describe the mechanis ms prior to the radiative relaxation, the dashed lines correspond to the s imulated population whe n the radiation recombination is turned off. It allows to highlight that the population evolutions of high order excitons considering radiative transitions (solid curves) is almost superimposed to the curves when the radiative recombinations are turned off (dashed curves) evidencing that the Auger quenching is dominating. We thus did not cons ider photon e miss ion of higher order exciton than XXX. In order to extract the trends for comparison with experimenta l data, the rat io of various emiss ion spec ies with the neutra l X are summarized in the Table 4.1.

Steady state radioluminescence

To assess the scintillation property, a series of SQWs were tested under X-rays illumination. As a first step toward probing the scintillation behavior in nanoparticles, we compared steady state emission s pectra under optica l and X-rays exc itations of 15nm, 25nm, 30nm and 40nm SQWs depos ited on s ilic on s ubstrate. Before depos ition, the solution of SQWs was diluted in a solution of hexane : octane (9:1). The mixture of hexane and octane can dry homoge neous ly avoiding the coffee -ring effect that concentrates the NCs on the border of the drop [140]. The films were obtained by dropcasting respective ly 40μL, 20μL, 20μL and 10μL of solution concentrated at 3.5μM, 0.87 μM, 0.57 μM and 0.77 μM and then the films were left to dry severa l minutes. Because the goal is to compare spectral shapes under X-rays and optica l exc itation, it requires to guarantee to be out of any s ignificant se lf-absorption regime (able to modify the spectral shapes), and radioluminescence is more sensitive to the thickness of the sample since the extreme ly strong penetrating ability of X-rays. Therefore, very thin films are required to perform the experiment. The steady state spectra were performed on a SR-500i imaging spectrometer from Oxford Instruments. When compared to the PL spectra of the solution sa mples (abs<0.1 at exc itation wavele ngth) as shown in Figure 4.8, no obvious shift and broadening were observed, indicating the self-absorption and energy transfer of the film sample s were negligib le . The accumulation time for photoluminescence and radioluminescence acquirements are set as 10s and 600s, respectively. The PL spectra were measured under a weak optical pumping at 405nm with a power of 3.5μW(unfocused). As show in Figure 4.9(blac k line), a s ingle band was observed, whic h is attributed to the s ingle exc iton (X) peaking at 1.914eV(648nm) , 1.908eV(650nm) , 1.907eV(650nm) , 1.906eV(651nm) for respective ly the 15, 25, 30 and 40nm particles. In the case of X-rays excitation (35kV, 35mA on W anode), the RL spectra exhibit a blue-shifted shoulder, and as the size increases, a small blue shift of the ma ximum of emiss ion is observed as the size increases (1.912eV, 1.912eV, 1.914eV and 1.917eV)(Figure 4.9(red line)). We attribute this blue-shifted shoulder indicated by the blue arrow to the multiexc itons emiss ions occurring under Xrays excitation contrary to weak intens ity optica l [START_REF] Turtos | Ultrafast emission from colloida l nanocrystals under pulsed X-ray excitation[END_REF]. The shift increases with the increasing shell, might be caused by the discrepancy of the exciton-exciton Coulomb interaction in different s ized particles [START_REF] Piryatinski | Effect of Quantum and Dielectric Confinement on the Exciton-Exciton Interaction Energy in Type II Core/Shell Semiconductor Nanocrystals[END_REF]. In addition, we also think the difference at low energy side can be partia lly re lated to trap-re la ted emiss io n in larger NCs [141]. Cons idering inte nse photoexc itation and X-rays excitation, traps and defects are fille d by a great amount of carriers, leading to a reduction of trap-re lated emiss ion. Therefore , Xrays excitation differs from low power optical excitat ion which cannot generate multiexcito ns. Since several studies have reported a rapid loss of emiss ion intens ity when the NCs was exposed under X-rays or γ-rays [START_REF] Withers | Rapid degradation of CdSe⁄ZnS colloidal quantum dots exposed to gamma irradiation[END_REF][START_REF] Gaur | Effects of x-ray and gamma-ray irradiation on the optical properties of quantum dots immobilize d in porous silic on SP IE Defense, Security, and Sensing[END_REF], it is necessary to examine the photostability of our samples under high energetic exc itation. Here, we applied the above sma llest and largest (15nm and 40nm) SQWs for experiments to verify the photostability of the samples. The experimenta l conditions are the same as above . The samples are placed about 5 cm in front of X-rays output w indow, and irra diated w ith X-rays(35kV,35mA) for up to one hour, during whic h RL s pectra of different irradiation times are acquire d, a nd all the accumulation times are set to 400s to facilitate the comparison of emission intensity changes at different irradiation times. As illustrated in Figure 4.10, we did not observe any emission intensity degradation with the increase in irradiation time, even the relative ly smaller 15nm SQWs show excellent photostability. This implies that the existence of thick-shell CdS achieves a well surface passivation, so that the sample is able to resist the photo-oxidation induced by ionizing radiation, thereby reduc ing the damage to the emission intensity. Since the accumulation time is 600s when the above steady-state RL measurements were performed, there is no need to consider the effect of degradation. And such a good photostability guarantees the reliability of the subsequent experimental results.

Photoexcitation

Photoluminescence decay curve measurement

Generally, steady state spectroscopy based on the ensemble samples does not enable the differentiation and observation of high order multie xc itons. In order to further figure out the specific excited state, we employ the time-resolved technique to observe the short-lived components and realize the spectra l and dynamica l ana lys is of multiexc itons under both pulsed laser excitation and X-rays excitation.

We first investigate the effect of shell thickness on the single exciton lifetime, by the mean of the decay curve measurement with 15nm-40nm SQWs. The excitation source is a Deltadiode from Horiba with the output light wavelength of 405nm and 100KHz repetition. The pulsed laser is defocused to avoid the formation of multiexc itons(the photons absorbed per NC on average <N><<1). The decay curves are shown in the Figure 4. 11(a). As the she ll grows , the lifetime of exc iton is continuous ly prolonged. The recombination rates strongly decrease in the larger NCs, which is consistent with the delocalization of the electrons into the shell, leading to a reduction of electron-hole wavefunction overlap [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF]. The PL decay should theoretically be well fitted by an exponentia l function. However, for large particles , apparently multie xponentia ls are needed to fit the decay curves. For instance, 40nm SQWs, the lifetimes we extracted are approximate ly 40ns, 180ns and a very long time constant 990ns. Here, we make a hypothetical ass ignme nt of these lifetimes. First, we assign the longest lifetime to the trap-related dela yed emiss ion. It involves the capture of the charge carriers of the traps , and then a time-de lay due to detrapping process, subsequently, the radiative recombination of the recovered carriers at 1S 3/2 1S e exciton state results in such a long lifetime.

Secondly, the assignment of time constant 180ns based on the reference [START_REF] Jeong | Colloida l Spherical Quantum Wells with Near-Unity Photoluminescence Quantum Yield and Suppressed Blinking[END_REF], in which the single exciton of largest SQWs they prepared (15nm in diameter) is characterized with around 60ns lifetime. Owing to significant reduction of the electronhole wavefunction overlap in our 40nm SQWs compared to the 15nm they studied, we can expect the increase in the single exciton, thus the time constant 180ns is assigned to single exciton. Due to the inevitable charging effect, the shorter time constant 40ns probably comes from the charged exciton. The fitted single exciton lifetimes of the ensemble samples are plotted in the Figure 4.11(b). Note that the single exciton of 15nm SQWs is around 35ns, which is a little smaller than the reported value 60ns. Compared to the one they reported with similar sizes (1.3/0.3/5.8nm), our samples (1.3/1.8/4.4nm) with a thicker CdSe emissive layer can lead to the increase in electron-hole overlap integral, which is responsible for the reduced lifetime. Additionally, we would like to point out dela yed emiss ion in large partic les extends the lifetime of s ingle exc itons to a certain extent, so the fitted value may be overestimated compared to small particles. In the follow ing section, we will focus on 15nm, 25nm and 40nm SQWs. To observe and ana lyze the dynamical evolution w ith varying exc itation intens ity, we first performed the decay curve measurement to record the decay kinetics of drop-cast 15 nm SQWs under different power photoe xcitation, varying from 1nW to 35 W, the diameter of the laser spot is around 4 m. In Figure 4.12, the decay curves were recorded at three different energy positions on the spectrum, which is illuminated by a low power (a)(<N>~0.003) and a relatively high-power photoexc itation(b)(<N >~4). For low power, the PL spectrum is only composed of single exciton, although the detected wavelengths are different, the decay curves obtained are consistent, indicating that they are derived from the same excitation species recombination. When <N> is around 4, which makes sure the possibility to generate the multiexcitons, we observe significantly different decay behavior, where the proportion of fast components gra dua lly increases with the blue s hift of the detection position. At 2.14eV, the rapid decay is on the order of the instrument response (IRF). This indicates that there are different recombination pathways at these three positions, involving the formation of single , charged and multiexc itons. Next, we will discuss the spectra and dynamics of these excitation states in detail. Since multiexc itons states can be generated at high pump powers by rapid successive absorption of multiple photons , the most common way to identify the excitation species is the analysis of their power-dependent properties. We focus on the dynamical signature at bandgap band(~1.9eV) to identify the species specifically. The decay curves show the strong power dependence, as illustrated in Figure 4.13. Specifica lly, at low power, the photons absorbed per NC on average <N>~0.003 ens ures that there is no formation of multiexc itons , emiss ion s hould be almost e xclus ive ly from the single exciton state. However, because the charging effect in NCs is usually unavoidable, even when the excitation optical intensity is very low, the entire decay curve still has a part of the contribution from charged excitons. As the excitation power increases, we normalized the PL decay traces at 60 ns to match the long-lived single exciton emission tail. We observe the emergence of additional fast components in the early time, corresponding to the multiexcitons. And the traces become indistinguishable at t >15 ns, indicating that all of the multiexcitons have already decayed and only single exciton left in the rest of recombination process. The method of extracting different exc ited states from the tota l time decay curve is based on the lifetime differences of these phys ica l events. By exploring the spectral distributions and exc itation power dependences, the excited states can be distinguis hed and assigned according to their unique spectral and temporal characteristics. Here, at lowest intensity excitation, the decay curve in Figure 4.13 exhibits two exponential dynamics by 12ns a nd 31. 7ns lifetime. Based on our ana lys is above, these time constants can be assigned to the charged exciton and single exciton emission. There may be a slight difference in lifetime due to the size distribution of the NCs. This is because the sample is placed on the focused spot in this measurement. At high excitation power, in order to quantify the lifetime of multie xcitons, we subtract the contribution of s ingle exc iton obta ine d by scaling the curve measured for <N> << 1, from the traces measured for high power, which allows us to isolate the purely fast components dynamics. As shown in Figure 4.14(a), we obtain the multiexcitons decay at the power of 1.6μW(<N>~4). The extracted dynamics can be fitted to a two-expone ntia ls decay with 0.72ns and 3.14ns lifetimes. We also found that with the increase of excitation power, these lifetime components gradually decrease.

Next, in order to validate the assignment of the fast decay component to multiexc itons recombination, we plot the relation of fast component integra l PL intens ity in the first ns with respect to the power fluence, as illustrated in Figure 4.14(b). The analysis of power-depende nt dynamics revea ls the compos ition of the superpos ition pe ak, that is a mixture of various excited states. First, as we expected, the slow component extracted at long-lived region(~100ns) w here the fast multiexc itons has a lready recombined scales linearly with the laser intensity. This behavior is in agreement with the single exciton feature. On the other hand we observed a nearly quadratic(~1. 88) dependence of the fast component integra l intens ity with respect to power, indicating the absorption of at least two photons w ithin the same pulse by a s ingle SQW. Thus on the basis of these measurements and analys is , the biexc iton and higher order(>2) multiexcitons are characterized by 3.14ns and 0.72ns, respectively. In order to elucidate the effect of shell thickness on the performance of multiexc itons, we further exa mine the power dependent PL dynamics at bandgap emiss ion band of 25 nm and 40 nm SQWs w ith the same e xcitation conditions , as shown in Figure 4.15. To match the single exciton emission tail, we normalized the decay curves of various exc itation fluences at 100ns a nd 250ns , whic h a llows us to highlight the emiss ion of multiexc itons. They show extremely different behavior in comparison with that of 15nm SQWs especially the 40nm sample. At low pump level, the decay curve of 40nm SQWs predominates by a significantly long-lived single exciton(~180ns). The delocalization degree of the electron spatially decreases the overlap between electron and hole wavefunction, whic h is respons ible for this low recombination rate. Besides , as the power enhanced, multiexc itons aga in e merge with cons iderably long lifetimes, the fast component is not completely decayed even at 100ns. Also, by fitting the decay curves at different powers, we have extracted the characteristic lifetimes of multiexcitons. As a summary, Table 4. 2 shows the four types of exc ited states w ith different lifetimes for 15nm, 25nm and 40nm SQWs, and further gives the hypothetical assignments of each type. The time constants are extracted at <N> ~4, 3 and 3 for 15nm, 25nm and 40nm SQW, respectively.

According to the statistical scaling described by our earlier study [START_REF] Mcguire | N ew Aspects of Carrier Multiplication in Semiconductor Nanocrystals[END_REF] , τ r , BX = τ r , X /4, using measured single exciton radiative lifetimes, we are able to estimate that the radiative lifetimes of biexciton, τ r,BX , are ~7.9ns , 18.5ns and 45ns , respective ly. Then we can calculate the qua ntum yie ld of biexc iton are around 40% , 23% and 22% , whic h are similar with the values reported for alloyed NCs recently [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF][START_REF] Bae | Controlled Alloying of the Core-Shell Interface in CdSe/CdS Quantum Dots for Suppression of Auger Recombination[END_REF]. Contrary to our expectation, the biexc iton quantum yie ld of large sized samples is relative ly low. Since the de layed e miss ion in large partic les causes the lifetime of s ingle exc iton to be overestimated as we mentioned in section 4.4. 1, we be lie ve that the quantum yie ld of biexciton in large particles is underestimated here. Considering that the biexciton lifetime is composed of radiative lifetime and non-radiative lifetime, = , + . (here we neglect non-radiative decay channels other than Auger recombination)[144] , for 15 nm SQWs, τ BX = 3.1ns and τ r , BX = 7.9ns, and we obtained τ A , BX of 5.1 ns. And we also estimated the Auger lifetime of biexc iton is 5.6ns and 12. 9ns for 25nm and 40nm SQWs , respectively. Obviously, compared to the typical Auger recombination lifetime in NCs, which is a few hundred picoseconds [START_REF] Klimov | Quantization of Multipartic le Auger Rates in Semiconductor Quantum Dots[END_REF], it is extended by nearly up to two orders of magnitude in our samples. Similarly, we can conc lude that the quantum yie lds of charged excitons are 76%, 58%, and 44% , respectively, and the corresponding Auger recombination lifetimes are 49ns, 50ns, and 72ns.

Time-resolved photoluminescence spectrum

In addition to the decay curve measurement to record the dynamics, in the thesis, time-resolved emission spectrum measurements were also employed to directly observe and analyze the changes in spectra with varying excitation intensity. First we performed the time-resolved PL spectrum measurement of drop-cast 15nm SQWs under the same optica l exc itation s ource(405nm) a nd intens ity range, from 9.3nW to 35μW. In order to investigate the ultrafast recombination of multiexcitons, here we can extract the PL spectra of multiexcitons through integrated spectra at different time intervals. As presented in 4.16(a) , the time-resolved PL spectrum were recorded when the incident laser is at a low power, approximate ly 9. 3nW, corresponding to the average number of excitons formed per NC per pulse <N>~0. 02. It is well known that multiexc itons in larger quas i-type II QDs induce the blue-shifted features (repulsive interaction) with respect to single exciton [START_REF] Nasilowski | Gradient CdSe/CdS Quantum Dots with Room Temperature Biexciton Unity Quantum Yield[END_REF]. Here, from the time gated spectra comparison of early time and late time, we cannot observe the signature of blue-shifted emission under low intensity excitation. The identical spectra manifest the absence of the multiexcitons emission and also indicate the homogeneous of this sample. If the sample is polydispersed, a red shift should be observed due to the energy transfer from sma ll to large partic les [141]. In addition, a lmost no de layed e miss ion was observed for 15nm SQWs, indicating that the sample is nearly trap-free, which is attributed to good surface passivation and release of interface strain.

When considering the high intensity optical excitation, the spectra evolution is significantly different. Here, we also performed the time-resolved PL spectrum measurements under higher power illumination with the same film sample. The power is around 35.78 μW, corresponding to <N>~90. The time gated spectra are illustrated in Figure 4. 16(b). Distinct emission features emerge in the early time at high energy side, and as the time increased, the features decay fast till disappear, leaving a single exciton spectral profile. These features are strongly power-depe ndent, which is the typica l signature of multiexcitons. In order to further interpret these distinct features, we plot the comparison of first 800ps time gated spectra under different inte ns ity exc itation, as shown in Figure 4.17(a). Note that in genera l, the regular drops observed are due to a small oscillatin g jitter of unknown origin having a significant impact of the first channe ls. We are allowed to directly see the power dependence of the excitation species. The various emission features with significant different dynamics are observed in sequence as the excitation power progressive ly increases for 15nm sample, w hich are named as Peak A, B, C and D. According to the emission energy, the features can be divided into two types: core states emission and shell emission.

Here, on the basis of the lifetime fitting results from the decay curve measurements of 15 nm SQWs, we first give a spectral assignment to the features originating from core states. At bandgap peak (Peak A), compared to the s ingle exciton spectrum (1.906 eV, la te -time spectrum in Figure 4.17), there is around 6meV blue-shifted (1.912 eV). Generally, exciton-e xciton interaction energy is on the order of 10meV in biexc iton, thus we assign it to the superpos ition of charged exciton and biexc iton. In addition, a pronounced emiss ion peak at 1.98 eV (Peak B) and a third peak at around 2. 11 eV (this peak is clearly distinguishable in larger particles in the following section, Peak C) can be seen. Considering the Coulomb effect of multi-carriers, the peak with 80 meV blue-shift relative to s ingle exciton emiss ion is assigne d to 1S-1S trans ition involving states with more than four charge carriers [START_REF] Vanmaekelbergh | Shape -Dependent Multiexciton Emission and Whispering Gallery Modes in Supraparticles of CdSe/Multishell Quantum Dots[END_REF] , whic h has been proved the dominant pathway in mu lt ie xc ito ns [START_REF] Shulenberger | Multiexciton Lifetimes Reveal Triexciton Emission Pathway in CdSe Nanocrystals[END_REF], inc luding charged biexc iton, triexc iton and higher order. In addition, partia l multiexc itons can a lso be recombined by 1P-1P and 1P-1S trans ition [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF], according to the research of Efros, Al. L. on CdSe NCs electronic structure [START_REF] Al | The Electronic Structure of Semiconductor Nanocrystals[END_REF]. We deduced that the higher feature at around 2.11eV originates from the P-le ve l re la ted e miss io n. In comparison w ith 2nd derivative of abs orption spectrum in Figure 4.17(b) , the stokes shift of bandgap peak is around 50 meV, and the P-leve l peak stokes shift is als o around 50meV. Fina lly, another band (Peak D) at around 2.4 eV have also been observed under high inte ns ity exc itation. Ow ing to pos ition of Peak D closes to CdS ba ndgap energy, here we assign it to shell-related optical transition [START_REF] Bae | Controlled Alloying of the Core-Shell Interface in CdSe/CdS Quantum Dots for Suppression of Auger Recombination[END_REF]. When a large number of excitons presence in the nanoparticles, the electron-hole pairs recombine directly in shell states to emit high energy photons instead of thermal relaxation to core energy states. We also carried out the time-resolved PL spectrum with 25nm and 40nm SQWs under the same laser power range. The time gated spectra of the highest power exc itation are shown in Figure 4.18 (a, c) : : we can also observe the four distinct emiss ion bands , originating from multiexcitons recombination. The feature shift with respect to the single exciton is summarized in the Figure 4.18(b). Compared to that of 15nm SQWs, the blue shift of each peak increases with the shell increasing. Peak A in 15nm SQWs is almost at the same spectral position as the single exciton peak (only 6meV shift). However, for 25 nm and 40 nm SQWs, this shift is approximate ly 20 meV and 30 meV, respective ly. The shift of Peak B is around 95 meV and 102 meV, and Peak C is blue shifted by around 240 meV and 260 meV. The reason for the trend of the higher energy in comparis on to the single exciton for the largest SQWs is that the delocalization of the electrons over the shell region increases with an increase in the shell thickness, this results in a more type II like behavior companied by the increase in the exciton-e xciton repuls ive interaction [START_REF] Piryatinski | Effect of Quantum and Dielectric Confinement on the Exciton-Exciton Interaction Energy in Type II Core/Shell Semiconductor Nanocrystals[END_REF][START_REF] Cihan | Observation of Biexcitons in Nanocrystal Solids in the Presence of Photocharging[END_REF]. This measurement is very sensitive to the focusing, slight deviation will make a strong effect on the ratio of the excitonic species, not on the peak position.

Besides, significant long-lived emission can be seen from the time gated spectra with a slight broa dening on the red side of the spectra at late time in Figure 4. 18 (d), this delaye d emiss ion has been demonstrated as the trap-related emission. Since the spectrum of the de layed emiss ion is s imilar to the single exc iton, it can be expla ined that de layed emiss ion is the result of recovery of trapped carriers and emiss ion from 1S 3/2 1S e exciton state. Usually, the delayed emission peak position exists slight redshift compared to the typical exciton, indicating the possibility of carrier trapping in smaller NCs is less likely than in the larger ones which might have more traps [141]. 19 presents the comparison of total emitted photons between low(black line) and high(red line) power excitation, which are consistent with the steady state PL. As we expected, the spectral profile under high power excitation is analogue to that under Xrays excitation, that is , a shoulder is observed on the blue edge , involving the formation of multiexcitons and, as the shell thickness increases, the spectrum under intense photoexc itation shows a slight blue shift around 3 meV, 6 meV and 12 meV for 15 nm, 25 nm and 40 nm SQWs, respectively. These observations reveal the energy recombination under X-rays radiation might undergo the similar process with that of inte nse photoexc itation. Meanwhile , it a lso points out that to compare the photoexc itation is a feasible way to investigate the excited state in radioluminescence. Therefore, in the following section we will further give the decay curve and time-resolved RL spectrum under X-rays excitation, and then explore the composition of its emitted light through comparative analysis.

X-rays excitation

In this sub-section, in order to investigate the exc itonic behavior under ionizing radiation, similarly to optica l excitation we carried out the time-resolved measurements with X-rays excitation (acceleration voltage 30kV on W anode) Through the comparison with the photoexc itation, we qualitative ly analyzed the dynamica l and spectral properties of excited states under X-rays excitation.

First, as shown in Figure 4.20, we recorded the decay curves at three different emiss ion wave lengths of 15nm SQWs. To acquire the s ignal w ith a good signa l to noise ratio, the integration time is set to 3600s because of the weak signal due to the low X-rays energy absorption by the very thin layer of partic les. The multi-exponentia l curve indicates the emission might be composed of several different excitonic species or experience quenching. Besides, a drastic change in the time response can be also observed under X-ra ys excitation as a function of the emission wave length. Cons istent with intense optical excitation, as the wavelength decreases, the curve decays more and more rapidly. This reveals that different types of recombination processes might occur at short-wavelength positions, which might be the signature of high order excitons presence.

Time-resolved radioluminescence spectrum

Separating emission both s pectrally and temporally allows us to observe multiexcitons features. The same strategy is thus used to highlight and compare X-rays induced multiexc itons emiss ions at short time. First, tempora l evolution of the emiss ion spectra of the samples after the pulsed X-ra ys excitation is given in the Figure 4.21. This measurement is set to 400s/point in order to obta in an acceptable signa l to noise ratio. And the s lits of the monochromator are ope ned four times larger(1mm) than that of photoexc itation (0.25nm). As can be seen in the Figure 4.21, it shows similar trends to intense optical excitation for all the particles. This observation is cons istent with the response of catholuminescence in nanocrystals reported by Lazaro A. Padilha [START_REF] Padilha | Response of Semiconductor Nanocrystals to Extremely Energetic Excitation[END_REF]147].

Because of the weak signa l, we use here 0-2ns time gate width ins tead of 0-0.8ns. Although the different emissive features are not as easy to distinguish as that of the 0-0.8ns interva l, we still observed a significant spectral broadening on the blue edge at early time. The emission peak at 2.0eV and 2.4eV can still be clearly identified, and another weak emiss ion band is between 2.1-2.2eV. Then the blue-shifted peaks gradually disappear as time passes after the excitation pulse hits the sample, indicating their fast decay dyna mics. At longer time gates, the spectrum shifts towards longer wave lengths by different energy over time. Similar to the photoexc itation, we compare the shift of the bandgap peak w ith respect to s ingle e xciton pos ition for three samples , whic h are approximately 6meV, 22meV and 27meV for 15nm, 25nm and 40nm, respectively. It is instructive to directly compare the RL and PL data for all sized particle. We compare the first 2ns emission under X-rays and intense optica l exc itation(35. 78μW) to highlight the fast component features as illustrated in Figure 4.22. We clearly see the similar spectral profiles under the two excitations, the difference just lies in the ratio of each excitation state. Based on this remarkable spectral similarity between PL and RL, we are able to assign these features to the multiexcitons under X-rays excitation, including charged exciton and biexc iton emitte d near the bandgap, and the ultrafast emiss ive bands raised from S-leve l trans ition and P-leve l trans ition involving the multicarriers (>4). These observations imply that as we predicted, during the se condary electron energy relaxation process, the electron-electron elastic scattering, carrier multiplication and Auger ionization pla y an importa nt role, promoting the formation of considerable electron-hole pairs and making SQWs occupied with high-order exc itation species.

Radioluminescence decay curve measurement

In order to further confirm the assignment of exc itonic spec ies, we employ the timeresolved technique to record the decay dynamics under X-rays excitation for three samples. Before this, for comparison, we first performed the decay curve measurement of the optical excitation in the non-focused condition to ensure that there is no formation of multiple excitons. At the same time, the non-focus ing s ituation is similar to X-rays irradiation (non-focused as well) regarding the particle size distribution, that can be then ignored.

Figure 4.23 illustrates the comparison decay dyna mics under X-rays and low power photoexc itation of 15nm, 25nm and 40nm SQWs. Both of the curves are measured without wavelength selection, including all the possible decay pathways for excited states. First, as we can see, the similar trend at relatively long time allows us to normalize the curves, we norma lized the curves at 60ns, 100ns and 200ns respective ly, and the results are analogous to the norma lization under different power excitation. The good match at the tail fully shows that single exciton recombination is the final state of the whole recombination process. Second, similar to high power light excitation, the obvious extra fast components can be observed under X-rays exc itation, and the multi-exponentia l curve reveals that there is not only one type of exciton. In addition, the fast component exhibits a strong size dependence in terms of decay time. As the particle size increases, the decay rate becomes significantly slower. For 15nm SQWs, the fast components completely decay after 20ns, however in 40nm SQWs, until 150ns there are still a part of emitted photons that do not originate from single excitons.

Next, we try to fit the decay dynamics by multie xponentia l in order to extract the fraction of extra fast signal from X-rays excitation reflecting the multiexcitons generation. First, according to the fitting results we discussed above, by fixing the charged exciton and single exciton lifetime, the PL decay curves are well fitted, we also extracted an additional long-lived c omponent, de layed e miss ion. Afterwards, by fixing single exciton lifetime or subtracting the contribution of s ingle exciton from the decay curve measured for X-rays, we are allowed to extract dynamics of the fast components. The fitting curves are shown in the Figure 4.23(red curves). The specific fitted parameters are given in Table 4.3. Here, we need to point out that the fastest parameter reaches the response function of the instrume nt. For 15nm sample , the fast components dynamics of RL decay curve are characterized by the 0.6ns, 2.6ns and 8.5ns lifetimes. These time constants are close to lifetime signatures extracted under intense pulsed laser excitation in Table 4.2 w ith a s light decrease. As we mentioned above , the extracted multiexc itons lifetimes gradually decrease with the increasing power, so this can be the interpretation for the s light discrepancy. Based on our ana lys is, we are able to ass ign the time constants in the Table 4.3, in where and are single and charged exciton, a nd then , are the lifetimes of multiexcitons : biexciton and higher order. Fina lly, we discuss the ratio of excite d states in different sized samples under X-rays excitation. In order to compare with the theoretical simulation, here we will focus on analyzing the ratio of each excitation species to single exciton, / , where the represents the amplitude of the excitonic species, is the lifetime. Considering that there are delayed emission photons in the emission spectrum, especially for large particles, which is ignored in the model, here we replace the single exciton with the sum of the single exc iton and the de layed emiss ion photon. Because the delayed emiss ion ma inly comes from the trapping and detrapping process of carriers, and the carriers trapping before the occurrence of carrier multiplication can result in a decrease in the number of multiexc itons. Therefore, we estimate that the ratio of multie xcitons we obta ined should be smaller than the simulation. 4.4 shows the specific fraction of different excitonic species to single exciton in varying sized SQWs. First, compared with the simulated values in Table 4.1, there is a quite significant difference in the ratio of charged exciton, which becomes dominant in compar iso n w it h mu lt ie xc ito ns. The simu la t io n d oes n ot take t he part ic le charg ing due to defects into account. However, the actual synthes ized NCs always present a certain number of defects on the surface or within the particles, which is unavoidable. Moreover, due to the more defect states in the large-s ize d partic les, this discrepancy is more significant, which is cons istent with the large amount of de layed emiss ion in the 40nm SQWs. For biexciton, with the increase in particle size, the ratio increases slightly, but not obvious ly, which shows the conformity with the simulation results. We belie ve that the absence of size dependence is due to the fact that the average deposited energy of each particle is quite high after X-rays excitation, resulting in very little influence caused by s ize, a nd that Auger que nching becomes the domina nt process . Since the lifetime of the higher order excitons is limited by the response function of the instrument, we will not discuss it in detail in this thesis.

Conclusion

In this chapter, the scintillation response of SQWs under X-rays excitation is analyzed using steady-state spectroscopy and time resolved technique. We mainly achieved the identification and quantitation of the exc ited stated under X-rays exc itation by comparing tempora lly and spectra lly the recombination produced follow ing high intense photoexcitation with X-rays ionizing radiation.

First, according to the unique temporal and spectral characteristics of different excited states, we realized the assignment of the excitonic species for intense optical excitation. Based on the extracted characteristic lifetime of each excited state, we calculated the quantum yie ld of bie xciton up to 40%, indicating that SQWs have excellent performance in suppressing Auger recombination. The remarkable spectral similarity between PL and RL enables us to analyze the distribution of the RL spectrum. Combined with the analogous decay dynamics, we demonstrated that under X-rays excitation, all observable emission can be attributed to recombination between states within the quantum-confine d NCs, including single exc iton, charged exc iton a nd multiexcitons. This conclusion reveals that carrier recombination plays an important role in the relaxation process. Secondly, we simulated the evolutions in the proportion of each excited state with respect to the size of the particles under X-rays radiation. The comparison between the experimental and theoretical value consistently shows that under the excitation of high-e nergy rays, the proportion of exc ited states has no obvious s ize dependence, as observed in the decay time analysis and time resolved emission spectra, which may be the result of the large amount of energy depos ited per NC on average. We suggest that Auger quenching is dominating when high order multiexc itons are produced in the confined structure which are in principle favored for larger sizes.

Chapter 5

Zhu MENG 107 Chapter 5

Scintillation properties of hererostructural II-VI nanoplatelets under X-rays excitation

With the ir appea ling properties , colloidal nanoplate lets have become highly promising for many applications. Their ultrafast exciton recombination rate caused by the giant oscillator strength transition received a lot of attention on the radiation detection fie ld [START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF][START_REF] Turtos | Ultrafast emission from colloida l nanocrystals under pulsed X-ray excitation[END_REF] , together w ith large absorption cross section and suppressed Auger recombination of heterostructures make 2D NP Ls potentia l scintilla ting materials.

Similar to SQWs, in this chapter we aim to describe the scintillation properties of NPLs under X-rays excitation via steady state spectroscopy and time resolved technique. We also compare the behaviors of exc ited states under inte nse laser excitation spectrally and dynamically to analyze the scintillation performance of different heterostructura l NPLs as a function of the morphologic aspects, including core/crown, core/shell and core/crown/shell.

Sample preparation and stability

The first step is to obtain efficient luminescence in NPLs by optimizing the quantum yield. The extended lateral area increases the interaction between NPLs as they form into a film, resulting in the stronger Forster resonance energy transfer (FRET) compared to the conventional QDs [START_REF] Guzelturk | Stacking in Colloida l Nanoplatelets: Tuning Excitonic Properties[END_REF]. Assisted by FRET, the possibility of charged carrier trapped by the defects considerably enhances, leading to rapid non-radiative recombination and significant decrease in photoluminescence quantum yie ld. In order to obta in bright film samples, we thus need to minimize the influence of the FRET on the optical properties. Owing to FRET distance sensitivity, to increase the spatial separation of the NPLs core should be the efficient method to avoid the FRET. Here, we take the core/shell NPLs as an example and prepare a NPLs-in-polymer sample. Specifically, the 5 mL NPLs was precipitated and redis persed in 1mL toluene which is benefic ial to the diss olution of PMAO (poly male ic anhydride-alt-octadecene), 50 mg PMAO powder was added to keep the volume fraction of NPLs less than 2% in the polymer matrix [START_REF] Guzelturk | Stacking in Colloida l Nanoplatelets: Tuning Excitonic Properties[END_REF]. Then the solution was placed on a silicon substrate via the drop-casting method. To exclude the selfreabsorption effect, only one drop(12 μL) was used to form an around one square centimeter film.

For comparison, we also prepared a core/shell NPLs film with the same quantity of particles but without the mixing of the polymer. The PL spectra of two film samples were acquire d in the same condition: illuminated by 405 nm low intensity diode and accumulated for 10 s. As illustrated in Figure 5.1(a), we observed more than 50% PL intensity decrease for the pure core/shell NPLs compared to the one embedded in polymer. Furthermore, a s light red shift (around 1nm) can be seen in the pure core/she ll NPLs spectrum, which is s imilar w ith the FRET s ignature of conventiona l QDs [START_REF] Crooker | Spectrally Resolved Dynamics of Energy Transfer in Quantum-Dot Assemblies: Towards Engineered Energy Flows in Artificial Materials[END_REF]. It can be interpreted by the inhomogeneous s ize broadening, favoring the energy transfer toward the partic les w ith s ma ller band gap. Since the inhomoge neous broade ning is induced by the shell growth for NPLs, the FRET of core NPLs with near-perfect monodispersity is usually not accompanied by red shift of PL spectrum [START_REF] Rowland | Picosecond energy transfer and multiexciton transfer outpaces Auger recombination in binary CdSe nanoplatelet solids[END_REF].

We also compared the time resolved decay curves of two film samples, as shown in Figure 5.1(b). In order to avoid the generation of multie xcitons, this measurement was performed under a very low power pulsed laser with 405nm exc itation wave length. We fixed the detected wavelength at maximum of PL spectrum. We observed that with the addition of the polymer, the recombination rate became slower, suggesting the fraction of the rapid non-radiative recombination decreased. Therefore, along with the enhancement of the PL intens ity, we conc luded the presence of polymer s patially separates the NPLs particles, and then effectively impedes the FRET between them. Similar to SQWs, we also tested the stability of NPLs under X-rays exc itation (35 kV, 35 mA, sample at 5 cm). Here, we continuously illuminated the core/shell NPLs (thickness ~4 nm) for around 2 hours and recorded the RL spectra at different time s, 15 min, 30 min and 100 min, as presented in Figure 5.2. Under long-term ionizing radiation, unlike SQWs, NPLs exhibits a certain degree of instability. We observed a gradual decrease in the RL intensity, wherein only 9% reduction can be seen for 30 min radiation, up to around 19% reduction for 100 min radiation. In this chapter, for the steady state measurements , we set the accumulation time as 400s which is much shorter than 30 min, thus the influence of RL spectra degradation should be very small. For time resolved RL measurements, the X-rays intens ity is generated by 30 kV voltage and 0.3 μA current, wh ic h is aro un d 1 0 5 times lower than that in steady state measurement. The acquisition time is usua lly set as one day whic h corresponds to ~9 s in the case of 35 mA current. Therefore, we also don't need to consider the degradation of luminescence intensity in time-resolved RL test. As they are promising fast time response radiation detection material, we will first discuss the scintillation behavior of core/crown NPLs. First, the PL spectrum was measured with 405nm diode (unfoc us ing, ~3.8μW). Owing to the only lateral direction crown coating, the PL line shape should be highly homogeneous , we thus fit the PL spectrum w ith Lorentzian function, as show n in Figure5. 3(a). However, the spectrum is not well fitted, indicating the presence of a slight size distribution. Then for comparison, Figure 5.3(b) plots the fitting curve with Gauss ian function, which is always related to spectrum cons isting of an inhomogeneous broadening. As we can see, due to the broadening induced by polydis persity is weak, the Ga ussia n function cannot be the best option for the fitting. Hence, in the following section, the spectra are fitted by the sum of the Lorentzian and Gaussian line shapes. Comparis on between PL and RL spectra is shown in Figure 5. 4(a, b). The profile in the left pane l shows a spectrally symmetric photoluminescence at low optica l excitation, which arises from the s ingle exciton emiss ion , and a broadening on the red edge with respect to the PL spectrum can be observed under 35 kV X-ra ys excitation (right pane l). The red shift has been previous ly reported under intense optica l exc itation, which is referred to the emission of biexciton [START_REF] Grim | Continuous-wave biexciton lasing at room temperature using solutionprocessed quantum wells[END_REF].

Here, to identify the excitation species with the red shift, we fit the PL and RL spectra using a pseudo-Voigt function, as depicted in following function, Where the second term is the s um of Lore ntzia n and Ga ussia n line s hapes with re lative contributions m u and 1-m u and widths and respectively. A is the amplitude.

As shown in 5.4(a), the PL spectrum at low power is well fitted by the single pseudo-Voigt function V 1 , the energy of symmetric peak is around 2.39eV with 47meV line width. In the Figure 5.4(b), we needed a sum of two pseudo-Voigt functions y = A 1 V 1 + A 2 V 2 to fit the RL spectrum. The parameters obtained from above V 1 fitting are held constant for RL spectrum fitting, and V 2 which represents the second emiss ion peak is fixed as the same line shape with V 1 . And the n the second peak energy is obta ined at 2.36eV, as shown in the green line of Figure 5.4(b). The difference between two peaks energy is the result of exciton-exc iton Coulomb interaction, about 30meV binding energy. For 2D nanocrystals , the ratio of the binding e nergy of the heavy hole biexc iton to that of the heavy hole exc iton is found to be 0. 2, and nearly indepe ndent of well w idth [START_REF] Birkedal | Binding of Quasi-Two-Dimensiona l Biexcitons[END_REF]. The exciton binding energy for 4ML NPLs is demonstrated to be approximate ly 210 meV [START_REF] Zelewski | Exciton Binding Energy of Two-Dimensional Highly Luminescent Colloida l Nanostructures Determined from Combined Optical and Photoacoustic Spectroscopies[END_REF]. Therefore, the oretical prediction of the biexc iton binding energy of our sample should be around 42 meV. Here, we can conclude that a biexc iton e miss ion under X-rays excitation is present. Furthermore, the slightly smaller red shift(~10meV) compared to the prediction indicates the occurrence of other types of multiexc itons emission. Because the Stokes shift of C/C NPLs is very small, it will lead to strong selfabsorption. The red-shifted multiexc itons band is benefic ia l to reduce the influence of self-absorption and conduce to practical application.

Time-resolved measurements

Figure 5.5 shows the normalized time-resolved RL spectra using different time intervals, early time (0-0.8 ns) and later time (4-6 ns) after the excitation pulse. The pulsed laser emits at 405 nm w ith 500kH z repetition rate, and the X-rays illumination at a voltage of 30 kV voltage leading to a 0.3μA current. Since the low signal for X-rays excitation, spectrum was taken with a 1mm spectrometer entrance slit to improve light collection, which leads to a broader spectrum compared to the steady state one shown in Figure 5.4(b). The accumulation time is around one day, corresponding to 400s/point. Although there is a lot of noise in the spectral line, from the figure we can clearly see that at late time, the emission peak of this time interval is obtained at ~2.39eV, which is well in correspondence w ith the s ingle exciton of low power photoexc itation. The early time gate(0-0.8s) is emitted at 2.37eV, a red shift of 20meV compared to the late time emiss ion, which is cons istent w ith signatures of multiexc itons emiss ion observed in steady states RL. And the multiexcitons have already decayed at 4ns. Using time resolved technique, we performed the decay dynamic measurement for core/crown NPLs under X-rays and low-intensity laser excitation, as depicted in Figure 5.6. The excitation source is the same as above. Here, the decay under photoexc itation was measured at unfocus ing pos ition (23μW/cm 2 ) in order to facilitate comparison with X-rays. First, tri-exponentia ls function is applied to fit the curve of photoexc itation. The multiexponentia l behavior has norma lly been observed in ensemble and individual NPLs at room temperature [START_REF] Tessier | Spectroscopy of Single CdSe Nanoplatelets[END_REF][START_REF] Olutas | Lateral Size-Dependent Spontaneous and Stimulated Emission Properties in Colloida l CdSe Nanoplatelets[END_REF]. This shows that the several possible recombination pathways might exist in the NPLs. The amplitude-averaged lifetime is often calculated as avg =( Σ × )/Σ , where is the lifetime and is the corresponding amplitude. The avg under photoexc itation is around 1. 56ns. The specific fitting parameters are given in the Table 5.1. For X-rays excitation, we observed a similar trend on the long tail of the decay curve. Through fixing the longest lifetime parameter, we also fit the curve with tri-exponentia ls , and the avg is calculated as 0.75ns. The faster decay is attributed to the formation of multiexcitons. The lifetime we obtained is longer than that in the literature [START_REF] Turtos | On the use of CdSe scintilla t ing nanoplatelets as time taggers for high-energy gamma detection Npj 2[END_REF], which should be the result of the trap related emiss ion in our NPLs. Table 5.1. Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/crown NPLs under laser and X-rays excitation. The average lifetimes: avg =(Σ × )/Σ , were calculated by using the respective amplitudes of each component. Due to the effective protection from the surrounding environment through the coating of the shell, the core/shell NPLs has demonstrated the improve ment of chemica l stability and quantum yield [START_REF] Altintas | Giant Alloyed Hot Injection Shells Enable Ultralow Optic al Gain Threshold in Colloida l Quantum Wells[END_REF]155]. On the other hand, although part of the quantum confinement is relaxed due to vertical growth, compared with spherical QDs, it exhibits a faster characteristic lifetime. At present, there are few analyses on the scintillation properties of the core/she ll NPLs structures. A full understanding of different types of NPLs is helpful to provide more selectivity for practical applications. Therefore, here we study the core/shell CdSe/CdZnS NPLs with different shell thic knesses. The samples are prepared by dropcasting the mixture of NPLs and polymers on the s ilicon substrate , the effect of the polymers is to reduce energy transfer among the particles. At the same time, the quantity of partic les dropped is required to be low to a void the influence of se lfabsorption on the spectral lines.

We first performed the steady-state PL and RL spectra measurements, as shown in Figure 5.7. We compared the normalized spectra of the three samples with different shell thickness named as core/shell_s, core/shell_ m and core/shell_ l for 4ML, 5ML and 12ML shell, respectively. The power of the photoexc itation is always kept very low and unfoc used to ensure that the resulting spectrum corresponds to the spectral profile of single excitons. As we can see, the three samples emit at 2.01eV, 1.96eV and 1.84eV, respectively. The red curves in the figure are the RL spectra, measured at an accelerating voltage of 35kV. Under X-rays excitation, the emission spectrum of core/shell NPLs shows the similar prof iles with s ingle exciton emiss ion spectrum of PL. However, as the thickness of the shell increases, we observed an emiss ion band on the blue side. It behaves on the oppos ite to the red multiexc itonic emiss ion band of core/crown NPLs. It shows that the interactions between the excitons of core/shell NPLs are different from that of core/crown configuration. In addition, this band shows a size dependent behavior, that is, the feature becomes distinct as the shell grows. In order to analyze the origin of this emission ba nd, we have conducted time-resolved research on the samples, presented in the next section.

Time-resolved measurements

Using time-resolved technique, we first focused on the characteristic lifetime of the three samples w ith different she ll thicknesses. The same as the above experimenta l conditions, low-intens ity photoexc itation and the sample is in a defocused pos ition, therefore, the influe nce of multiexc itons is exc luded. As shown in Figure 5.8, first, the decay curves of the samples are all multi-exponentia l, which is common in NPLs as mentioned above. Multiple complex channe ls of excitons, including single exc iton radiative recombination, trap recombination and non-radiative recombination are considered as the origin of this multiexponential behavior [START_REF] Tessier | Spectroscopy of Colloida l Semiconductor Core/Shell Nanoplatelets with High Quantum Yield[END_REF][START_REF] Kunneman | Nature and Decay Pathways of Photoexcited States in CdSe and CdSe/CdS Nanoplatelets[END_REF]. The core/shell structure sample we used has a reported quantum yie ld c lose to unity [START_REF] Altintas | Giant Alloyed Hot Injection Shells Enable Ultralow Optic al Gain Threshold in Colloida l Quantum Wells[END_REF], w hich contradicts the existence of a large number of non-radiative recombination pathways. Here, we think the reason is that we are studying the thin film sample, in which the reduced distance between particles promotes the energy transfer between them, and the proba bility of charged carriers being trapped by defects is greatly increased, which leads to non-radiative recombination. Although the addition of polymers can reduce the energy transfer effect to a certain extent, it cannot achieve 100% elimination. Secondly, based on the multiexponential fitting, we calculated the amplitude average lifetime, w hich are around 7. 6 ns , 10 ns to 33 ns , respective ly. We can observe that the lifetime gradually increases with the growth of the shell thickness, which reflects that the relaxation of the qua ntum confineme nt as the shell increases makes the reduction of the wave functions overla p of electrons and holes. Besides , the fraction of long-lived component A 3 3 is large in core/shell_l, suggesting the presence of an increase in trapinvolved de layed emiss ion, and the traps usually arise from the lattice misfit defects in thick shell nanocrystals. The specific fitting parameters are shown in Table 5.2. To provide further ins ight into the emiss ion kinetics of various sized core/she ll NPLs, we also measured the power dependence of the decay dynamics using time-resolved spectroscopy. This a llows us to analyze the multiexc itons dynamica l process in the core/shell structure. Here, we mainly focused on the decay curves near the maximum of emission peak, which are 2 eV, 1.94 eV and 1.82 eV respectively. The excitation laser source is 405 nm w ith 100kH z repetition, the power range is measured from 0.07 μW to 33 μW, and the size of the laser spot is approximate ly 4 μm in diameter. Figures 5.9(a-c) summarize the power dependence curves of three samples. The traces of PL decay under varying power are norma lized for a long time (~20 ns) to highlight the fast component. We can clearly see that all the decay curves are well matched in the long tail, which indicates that the recombination of same excited state species takes place at this time. Similar with the fitting results in defocusing case, three different time constants are needed to fit the curves at low power excitation, and the observed slight deviation is caused by the size distribution of particles.

In addition, under high-power excitation, we can clearly observe the emergence of fast compone nts, which decay rapidly. The fast components obta ined by multiexponentia l fitting are around 0.7ns, 0. 9ns and 1.8ns. The obta ined fast component lifetime is s lightly higher than the typica l Auger recombination lifetime(hundre ds of ps), suggesting the reduction of Auger effect owing to the alloye d shell [START_REF] Bae | Controlled Alloying of the Core-Shell Interface in CdSe/CdS Quantum Dots for Suppression of Auger Recombination[END_REF]. Additiona lly, with the power increases, the lifetime of the fast component decreases gradually, as summarized in Figure 5.9(d). In 2D nanocrystals, different with QDs, the exciton is delocalized in the partic les instead of be ing c onfined. and the rate of A uger non-radiative pathways highly de pends on the initia l dens ity of the exc itons generated after laser excitation or the inter-exciton distance. For intermediate power, spatially separated excitons have to diffuse before encountering and recombination. The increasing power induces the increase in the density of excitons, leading to the decrease in the inter-exciton distance.The decrease in the diffusion time is responsible for the acceleration of the reconbination rate. When we consider high power excitation, the excitons can recombine before diffus ion, thus the recombination is us ually on the order of ten ps [START_REF] Franco | Etude par spectroscopies optiques résolue dans le temps des dynamiques ultrarapides dans les nanostructures colloïdales de pérovskite[END_REF]. However, time resolved technique is not able to analyze such a fast lifetime. We do not discuss it here. The specific fitting parameters are shown in Table 5.3-5.5. In order to further verify the origin of the fast component, taking core/she ll_ l as an example , we plotted the relationship between the PL intens ity of the slow and fast components with the excitation power, the slow component is extracted from the decay curve Figure 5.10(a) at 60 ns, where the fast component is fully decayed. And the fast component is isolated by subtracting the contribution of the single exciton from PL intens ity at 0 ns, as shown in Figure 5.10(b). We can see that the slope of the s low component as a function of power is about 0.7, which is less than the expected 1. Here we think that the reason is the intense laser has caused the degradation of the sample. Although the presence of polymer can alleviate a part of the degradation, it still induces a inevitable negative influence on the emission intensity. In addition, the fast component increases nonlinearly with the increase in power, which is the signature of multiexcitons emiss ion. H owever, the obta ined s lope va lue is only 1.3, greater than 1 but less tha n 2, which is the characteristic of biexcitons with power. The reason of the smaller slope might be the degradation. The evolution of PL intensity as a function of power: slow component is extracted from (a) at 60ns and fast component is the PL intensity at 0ns after subtracting the contribution of single exciton.

In order to investigate the types of exc ited states under X-rays radiation, we then compared the decay kinetic curves under photoexcitation and X-rays excitation. The curves are measured without wavelengths selection. The decay under photoexcitation is obta ine d in the case of defocusing (the same with Figure5. 8), and the laser power is kept very low to avoid the formation of multiexc itons. As presented in Figure 5. 11, we normalized the two decay curves at the long time: 20 ns, 25 ns and 40 ns respectively. As a result, all the decay curves under X-rays exc itation can well match the photoexc itation traces at the long tail, which indicates that these two different excitations produced the same type of long-lived exc ited state. In addition, the curves under X-rays excitation has an additional fast decay component, which is similar to the decay dyna mics under inte nse light excitation. Based on the similarity, we speculate that the fast component is caused by the formation of multiexcitons. In order to quantitative ly ana lyze the fast component under X-rays excitation, we subtracted the light excitation curves from the decay curves under X-rays excitation to obtain a pure fast component. Through fitting, we obtain that the fast components are characterized as 0.51ns , 0. 54ns and 0.64ns respective ly, and this time constant has reached the response function of the instrume nt. The specific fitting parameters are shown in Table 5.6, and the obta ined time constants a ll correspond to the fitting time under intense light exc itation in Table 5.3-5. 5. The slight de viation in the data is due to the size distribution of the samples. This consistent decay kinetic curve further demonstrates temporally that the s imilar types of exc ited state are generated under X-rays irradiation, name ly, multie xc itons through the carrier multiplication process. We can see that compared to core/crown NPLs, the average decay time is larger.

We calculate the fraction of the fast component us ing /(Σ × ), 24% , 26% , 13% for core/shell_s, core/shell_m and core/shell_l, respectively. Different with steady state spectra, no obvious size dependence of multiexc itons can be observed, especia lly for the core/shell_l. We think the decrease in the fraction might be attributed to the capture of carriers by the defects in larger particles. For steady state measurement, the defects can keep being fille d by the continuous illumination of X-rays, the filling states result in the reduction of the carriers trapping and the long-lived components. Therefore, the possibility of the formation of multiexcitons is enhanced. However, in the case of pulsed X-rays, the long-lived emiss ion is 30% of the tota l emiss ion photons in the core/she ll_ l NPLs and almost negligible for core/shell_s and core/shell_ m, so part of the carriers can be captured before the occurrence of the carrier multiplication, leading to a loss of the multie xcito ns in core/shell_ l. As compared to low power optical excitation, we observe an acceleration of the decay under X-rays while the emission spectra do not change in small shell samples. We employed thus the time-resolved RL spectrum to investigate the multiexcitons spectra under X-rays excitation. Here we need to point out that in order to obta in high-quality spectra, the sample used for X-rays testing is increased in concentration in order to collect suffic ient signa ls. According to our discuss ion above , such processing will have an impact on the spectrum of X-rays excitation, and self-absorption may cause a certain red shift in the spectrum. This is consistent with what we observed in the experimental results, as shown in Figure 5.12. We compared the corresponding spectra of 0-2ns and long-live d. Although the existence of se lf-absorption will affect the spectral lines, we can still cle arly observe the evolution of spectral lines over time under X-rays excitation.

First of all, the right s ide of Figure 5.12 are the time gated spectrum corresponding to intense light excitation(33μW) for three core/shell NPLs. Contrary to previous reports regarding core/she ll NPLs [START_REF] Altintas | Giant Alloyed Hot Injection Shells Enable Ultralow Optic al Gain Threshold in Colloida l Quantum Wells[END_REF]159], the fast component we observed, that is , the feature of multiexc itons , appears at the blue edge relative to the long-lived s pectrum. We observed the fast compone nt spectral lines in 0-2 ns shows a blue shift, about 10 meV for all samples , not a red shift. This indicates the energy band a lignment in our NPLs might behave like quasi type-II, and the blue shift arises from the repuls ive interaction between the excitons (see Section 2.3.4).

Secondly, in addition to the blue shift of the spectral line, in core/shell_ m, we can also observe a weak emission band at ~2.04 eV, which is 60 meV blue-shifted relative to the main peak. Similarly, in core/shell_ l, we can also distinguish at the position of 60 meV blue shift (~1.92 eV), there is an emission ba nd. In addition, at 2.02 eV, another emission band is also observable. Compared to the main emission peak, this emission band has a blue s hift of about 160 meV. Since the shift of the bie xciton is us ua lly around 10meV, the two emiss ion bands may arise from higher-order multiexc itons emiss ion. NPLs with thick shells exhibit a larger absorption cross section, which might be the reason for the formation of higher-order multiple exc itons. Ow ing to the limitation of the experimenta l equipment, we cannot extract the corresponding characteristic lifetime s through time-resolved technique. Figure 5.12 als o plots the spectra of three samples under X-ra ys excitation. In order to facilitate comparison, we have selected the same time interval as that of intense light excitation. We found that a ll spectra for X-rays are broader compared to the spectra for photoexcitation, this is because they were taken with a 1mm spectrometer entrance slit to impr ove light collection, while 0.25 mm was used for optical exc itation measurement. First of all, similar to photoexc itation, we observed the emiss ion shoulder of the 0-2ns time gated s pectrum appears at the blue edge as observed in steady state experiments , but we did not observe the change in the spectral line on the red side , which ma y be caused by the mentioned broadening in the spectrum. As time increases, this emiss ion characteristic gradually disappears. In addition, interestingly, for core/shell_ l (Figure 5.12f), at 1.915 eV a weak emiss ion band can be observed. This is complete ly cons istent with the spectral characteristics of intense light excitation. Therefore, by comparison, we obtained the similar emission spectra, which allows us to prove spectrally that the type of excited state generated under X-rays excitation is cons istent w ith that under inte nse light excitation, that is , a recombination process of multie xcitons takes place during the energy relaxation of core/shell NPLs under X-rays excitation.

Through time-resolved decay curves and emission spectra comparison, we have realized the spectral and kinetic identification of excited states produced following Xrays excitation in core/shell NPLs. Combined with the steady-state spectroscopy, we highlight that under X-rays exc itations , the energy re laxation process of high-energy electrons in the core/shell NPLs is also dominated by carrier multiplication. This leads to the formation and recombination of multiexc itons , accompanied by the appearanc e of shoulders on the blue side. And as the size increases, the fraction of multiexc itons emission gradually increases. 5.4 Core/crown/shell CdSe/CdS/CdZnS nanoplatelets Core/crown/she ll(C/C/S) NPLs is a novel two dimens iona l heterostructure [START_REF] Kelestemur | Platelet-in-Box Colloidal Quantum Wells: CdSe/CdS@CdS Core/Crown@Shell Heteronanoplatelets[END_REF][START_REF] Shendre | Ultrahigh-efficiency aqueous flat nanocrystals of CdSe/CdS@Cd 1-x Zn x S colloida l core/crown@alloyed-shell quantum wells[END_REF][START_REF] Taghipour | Sub-single exciton optical gain threshold in colloidal semiconductor quantum wells with gradient alloy shelling[END_REF], and so far there is little research to analyze its properties. Here we will brie fly discuss the scintilla tor behavior of this structure.

Steady state radioluminescence

Similarly, we first compared the normalized steady-state spectra of PL and RL of core/crown/she ll CdSe/CdS/CdZnS NPLs, as shown in Figure 5.13. Under low power photoexc itation, the partic le emits at 1.92eV. While for X-rays excitation, we observed that the RL spectrum has a slight blue shift relative to the PL spectrum, around 5meV. Compared with core/crown NPLs, RL does not show a redshift, which indicates that the existence of the 2ML shell changes the Coulomb interaction between excitons in the partic le , so its band a lignme nt shows a trans ition from type I to type II or quasi-type II [START_REF] Taghipour | Sub-single exciton optical gain threshold in colloidal semiconductor quantum wells with gradient alloy shelling[END_REF]. In addition, compared with the core/shell_s results, in which only the spectrum broaden toward the blue edge can be seen, the slight blue shift might reveal the increase in the repulsion between excitons or the different excited states are generated in core/crown/shell NPLs under X-rays excitation. Next, we conducted the experiments on the power dependence of steady state spectroscopy us ing solution samples. Since the quartz cuvette will produce scintillation signa l under X-rays radiation (even if the yield is weak, the fraction of X-rays photons in the quartz w ill be significantly higher), and meanwhile reduce the photons absorbed in the partic le s, t he solut ion samp le is not su itab le for X-rays related research. Therefore, this experiment is only used as an auxiliary research. The specific experimental device optica l parametric osc illation laser (OPO) and the preparation of the solution sample are described in Annex2. The partic le-dispersed solution can eliminate the influence of energy transfer between particles on the spectrum, and the preparation of lowconcentration samples can also reduce self-absorption. The normalized spectra are shown in Figure 5.14, as the excitation power continues to increase, the spectrum is broadened and accompanied by a slight blue shift, which is consistent with the performance of the RL spectrum. This directly indicates that similar radiative recombination may exist under two different excitations, and the effect of energy transfer between partic les in the film sample on the spectrum is negligible. The re levant spectra of all other samples can be found in the Annex2.

Time-resolved measurements

The PL spectrally resolved measurement was implemented at two powers, 33 μW and 4nW, as shown in Figure 5.15 for spectra at different time interva ls. First of all, in the case of low power, the 0-2 ns and 50-60 ns spectra are similar, which means that there is no additiona l fast component ge neration, and Forster resonance energy transfer is weak. Otherwise, the small-sized sample will emit photons after absorbing the emitted photons from large sized particles and then emits a photon with a lower energy, which will cause a red shift in the spectrum over time , regardless of the excitation power. In Figure 5.15(a), we did not find a distinguishable spectral redshift, so FRET can be exclude d, which is related to the good size distribution of the sample.

Next, we analyze the case of high-power exc itation (Figure 5. 15b, d). Here we divide the spectrum into 0-60 ns and after 60 ns and discuss separately. First, in the first 60 ns , we observed that the main emission peak at 0-2 ns is 3 4 meV b lue s h ifte d re la t ive to that at 50-60 ns. This s pectrally dependent emiss ion band is in agreement w ith the characteristics of multiexcitons. The blue shift is 3 times larger than that in core/shell NPLs with the similar shell thickness. This may be because the CdS crown promotes the delocalization of electrons in the conduction band, thereby enhanc ing the local charge imbalance and generating stronger repulsion between excitons.

With the increase in time, the blue shift of the main peak gradually disappears, and the broadening of the spectrum also gradually decreases. For details, shown in Figure 5.15(c), which describes the shift of the main peak of the 0-500 ns s pectra and the change of the spectral width. The delaye d PL spectrum after 60 ns is shown in Figure 5. 15( d). The shift of the main peak can no longer be observed, that is , it is similar to the emission of a single exciton, but it is characterized by a long lifetime. From this simila r it y, we c an c onc lu de that de la yed e miss io n is the resu lt of e miss io n fro m the recovered carriers at 1S 3/2 1S e exciton state [141]. In addition, in this time range, there is a slight broadening of the spectrum at the red edge, which may be due to the direct radiative recombination from the carriers trapped. Figure 5.16 is the norma lize d time resolve d RL spectra under X-rays excitation. Due to the low signal-to-noise ratio, the s igna l after 20ns is hidde n in the background a nd cannot be extracted. So here we can only discuss the first 20ns of the spectrum. We can see from the figure that, unlike the RL spectrum of core/shell NPLs, the early spectrum has a blue shift relative to the 10-20ns spectrum, whic h is about 33meV. Although the 10-20ns spectrum may still cons ist of a certain fast component , it can be known from the photoexcitation spectrum that there is only a slight deviation from the 50-60ns spectrum. Therefore, we can conclude that at 0-2ns the RL and PL under inte nse light exc itation exhibit a similar blue shift, and similar spectral broadening. This high consistency indicates that the recombination of fast components originates from similar exciton states with photoexcitation. We tested the decay dynamics curve at different powers, and the detection position is the maximum value of the spectrum, as shown in Figure 5.17(a). First of all, compared with other structures of NPLs, the C/C/S samples we synthes ized have a longer decay time. As the power increases, an additional fast component appears in the decay curve. We analyzed the evolution of the fast component and the s low component as a function of the excitation laser power in Figure 5.17(b). The slow compone nt is extracted at 75ns of the decay curve. The fast component is the 0ns PL intensity after subtracting the contribution of the single exciton. We can see that the slope of the s low component is close to unity, indicating that it changes linearly with the power. The slightly smaller slope may be caused by the degradation of the sample. For the fast component, the PL intensity exhibits a super-linear increase as the power increases. But similar with the result of core/shell NPLs, the slope is smaller than the simplest biexciton feature. The reason for this small slope is not clear, and sample degradation may be one of them. Nevertheless, this super-linear change represents that multiexcitons are formed in this process. Finally, we compared the decay kinetics curves under photoexc itation and X-rays excitation. Similarly, in order to facilitate the comparison, the photoexcitation curve is measured in the out-of-focus condition, this also ensures that no multiexcitons are generated. As shown in Figure 5.18(a), the two different excitation curves can be completely matched after 30ns, so we can subtract the light excitation curve from the Xray decay curve to obtain the pure fast component. The result is shown in Table 5.7 Summary of the fitting parameters for multi-exponential decay, lifetimes with corresponding amplitudes for core/crown/shell NPLs under low intensity photoexcitation and Xrays excitation, and the calculated average lifetimes. From the similarity on the spectra and the dynamics , we proved the formation of the multiexc itons under X-rays excitation. However, similar with the core/shell_ l, the C/C/S NPLs also suffer from the issue of delayed emiss ion, and the fraction of the de layed emiss ion is around 40%, which is cons istence with the low quantum yie ld. The fast component fraction is calculated as 10%, which is 2 times lower than that in core/shell_s with the similar she ll thickness. From this , we can know the defects strongly affect the quantity of multie xcitons through trapping the carriers before the occurrence of the carrier multiplication.

Conclusion

In this chapter, we apply steady-state spectroscopy and time resolved technique to analyze the scintillating properties of three different heterostructural NPLs under X-rays radiation.

2D NPLs have ultra-fast decay time because of their giant oscillator strength. Since the crown only covered the lateral direction in core/crown structure, it does not release the quantum confinement effect, and thus has a small influence on the decay dynamics curve. We obtain an average lifetime less than 1ns under X-rays excitation. In addition, compared to the PL spectrum, the RL spectrum exhibits a red shift due to the generation of multiexcitons , which is beneficial to reduce self-absorption and practical applications.

We also proved that the formation of multiexcitons under the radiation of X-rays also occurs in core/s hell and core/crown/s he ll NPLs based on the highly uniformity in spectra and kinetics compared to the intense optical excitation, indicating that the carrier multiplication phenomenon is effective in the 2D nanocrystals. But, unlike core/crown, the feature of multiexc itons a ppears on the blue edge. Meanwhile , due to the reduction of confinement and the delocalization of carriers, the decay time of the excited state in these two structures is longer, but it is still faster than most of QDs. Compared to the core/crown structure, they exhibit the better chemical and physical stabilities. For applications that do not require ultra-fast time response, they will have better performance. In addition, by analyzing steady state spectra of core/shell structures with different shell thicknesses, we found that in the partic le s ize range we are studying (less than 10nm in the vertical direction), the scintillator properties of the core/shell structure show a size dependent behavior. As the thickness of the shell increases, the emission ratio of multiexc itons gradually increases. However, for the pulsed X-rays, the size dependence cannot be observed, the interpretation might be that the carriers are trapped by the defects before the occurrence of carrier multiplication in large particles, leading to the decrease in the formation of multiexcitons.

Chapter 6

Zhu MENG 131 Chapter 6

Conclusion

This thesis ma inly focuses on the investigation of the sc intillator properties of II-VI group heterostructura l na nocrysta ls under X-rays ionizing radiation. We have completed the colloidal synthesis and characterization of nanocrystals and analyzed spectrally and tempora lly the optical response under optical and ionizing radiation aiming to describe the energy relaxation mechanisms and the types of excited states generated under X-rays excitation.

In particular, we successfully s ynthes ized s pherica l quantum we lls (SQWs) CdS/CdSe/CdS w ith varying she ll thicknesses, the dia meter is from 15 nm up to 40 nm with a good size dispersion. The sandwich-like structure enables the synthesis of giant shells without the loss of optical quality due to less misfit defects, and the quantum yield here is up to 80% for smaller particles, and even for the larger particles the quantum yield is measured as 35% in thin film. In addition, different heterostructural nanoplatelets (NPLs), core/crown, core/shell and core/crown/shell NPLs are prepared by hot injection method. We achieve d the formation of core/crown NPLs with regular sha pe, fac ilitating further overcoating and applications. And alloye d core/shell CdSe/CdZnS NPLs were fabricated in three varying she ll thickness with narrow PL linew idths, homoge neous morphology and high quantum yield. We also for the first time successfully coated a CdZnS shell on the core/crown seeds via the hot precursor injection with 20% quantum yield in thin film, demonstrating the possibility to fabricate core/crown/shell NPLs through hot-injection method.

To investigate the scintillation properties of NCs, we mainly analyzed the spectral and dynamical behavior under X-rays excitation via the steady-state radioluminescence technique and time-resolved techniques. We performed the decay curves measurement under varying excitation powers, this power dependence enables the identification of the excited states. Together with time-resolved emission spectra, we successfully achieved the assignment of single exciton, biexciton, and S-le ve l trans it io n an d P -le ve l trans it ion involving the multi-carriers (>4). The biexciton quantum yield was calculated up to 40%. Through comparative analysis between intense optical excitation and X-rays excitation, analogous emission features have been observed, which is an indication of the presence of similar excitonic species, proving that carrier multiplication plays an important role in the scintillation processes. Based on the analysis of decay dynamics, we studied the evolution of the ratio of each excitonic species in varying sized SQWs during the scintilla tion process. A simulation based on rate equations from Prof A.Vasil'ev (Moscow state univers ity) mode led the distribution of the excitation species production in different sized NCs under 30keV photon exc itation. The experimenta l results are consistent with the theoretical modeling, the fraction of the biexciton does not evolve significantly with the partic le size in the range we used, suggesting the excitonic species formed under extremely energetic electron excitation is insens itive to the nanopartic le size. We suggest that Auger que nching is dominating when high order multiexc itons are produced in the confined structure which are in principle favored for larger sizes. For this part, considering that larger size increases the decay time of the single exciton, we do not have benefit to use large shells for fast timing applications.

Compared to the SQWs, NPLs displayed faster lifetime due to the strong oscillator strength. Therefore, NPLs is an outstanding fast-timing scintillation material. By comparing the temporally and spectrally radiative recombination produced following excitation under photoe xciton and under X-rays excitation, we identify and quantify the distribution of exc itonic species during the scintillation process of different Cd-based nanoplatelets. Different behaviors have been observed in core/crown, core/she ll, and core/crown/she ll NPLs under X-rays excitation. Red-shifted spectrum with respect to PL was observed in core/crown and is characterized with fast time constant of less than 1 ns. This red s hift is benefic ial for the reduction of se lf-absorption. Meanwhile in core/shell and core/crown/she ll NPLs, the fast emiss ive features were found on the blue edge from the time resolved spectra. Based on the highly uniformity in spectra and kinetics compared to the intense optical excitation, we proved that the fast component origins from the multiexc itons under the radiation of X-rays, indicating that the carrier multiplication phenomenon is effective in the 2D nanocrystals. In addition, by analyzing steady state spectra of core/shell structures with different shell thicknesses, we found that in the partic le size range we are studying (less than 10nm in thickness), the fraction of multiexcitons emission gradually increases as the thickness of the shell increases. However, for the pulse d X-rays excitation, the s ize dependent behavior cannot be observed, the interpretation might be that the carriers are trapped by the defects before the occurrence of carrier multiplication in large particles, leading to the decrease in the formation of multiexcitons.

In summary, we demonstrated that the excitonic species during the scintillation process under X-rays excitation are similar to that under intense optical excitation in II-VI nanopartic les , both for SQWs and 2D NPLs. All observable emiss ion s igna ls can be attributed to recombination of single , charged and multie xc itons. The multiexc itons are mostly produced by carrier multiplication of secondary electrons. And the size-dependent trend is observed for sma ll s ized NPLs, not the larger SQWs partic les. These observations give direct ins ight into understanding of the energy de pos ition a nd relaxation process under X-rays radiation in nanoparticles.

Even though we have achieve d the investigation on the scintillation properties of II-VI nanocrystals, there are still many challenges that need to be overcome for future research and applications. First, nanocrysta ls whic h can suppress the Auger effect without the prolong of lifetime, are needed for fast timing scintillators. Meanwhile, since the defects usually introduce delayed emission, realization of defects free nanocrystals will also contribute to achieve this ultrafast decay. In addition, self-absorption limits the nanocrystals concentration, making the lower energy deposition in the scintillator, and also significantly affects the light collection emission of multiexcitons. Therefore, to reduce the self-absorption is also an important issue to address. Nevertheless, with the developing of the synthetic technique , the nanocrysta ls w ill be promis ing as a scintilla t in g materia l in the future.

Annexe1: Estimation of absorption cross section of spherical quantum wells

Zhu MENG 135 Annexe1: Estimation of absorption cross section of spherical quantum wells

We estimated the cross-section of the SQWs on the bas is of ref. [START_REF] Kholmicheva | Sustained Biexciton Populations in Nanoshell Quantum Dots[END_REF]. Specifically, the absorption cross-section of SQWs with size(1.3/0.6/5.3nm) at 3.1eV was measured as about 8.3× 10 -14 cm 2 by using the PL saturation method. We corrected the value with a factor of 1. 03 to take into account the exc itation wave length of our setup (3.06eV). Since the simila r size with o ur sma lle st part ic le s, we a re able to c o ns ider th is va lue is the absorption cross-section of 15 nm SQWs , which is 8× 10 -14 cm 2 . We are allowed to calculate the absorption cross-section of all the sized SQWs by assuming to the volume scaling [START_REF] Wood | Air-Stable Operation of Transparent, Colloida l Quantum Dot Based LEDs with a Unipolar Device Architecture[END_REF], as depicted in Table A1.1. The absorption cross section of the largest SQWs is around three orders of magnitude larger than that of regular quantum dots.

It should be pointed out that since the absorption at the excitation wavelength is mainly caused by the CdS shell, and the shell thic kness of our 15 nm samples(1.3/1.8/4.4nm) is a little smaller than the mentioned SQWs(1. 3/0. 6/5.3 nm) with the similar volume. Therefore this estimated value may be a little overestimated. To calculate the average number of exciton formed per nanocrystal per pulse(<N>), the following expression was used, <N>=j (A1-1)

where j and respectively represent the intensity of the excitation laser and the absorption cross section of the sample at the excitation wavelength. The excitation intensity of the laser can be obtained by, where the laser power P can be directly measured, f is the repetition frequency of the laser, A is the area of size spot, and, the λ, h, c are the excitation wavelength, Planck's constant, and the speed of light, respectively. The diameter of laser spot is approximately 4μm. Table A1.2 shows the <N> for different sized SQWs. Annexe2: Solution sample under optical parametric oscillation laser

In order to obtain steady-state multiexcitons emission characteristics, we carried out the power dependent PL spectra measurement with an optical parametric osc illation laser(OPO). Since too high laser intens ity is very easy to cause the film sample to be burnt out, thereby affecting the reliability of the experimental data, so for this measurement, the emission spectra have been performed in solution sample. All the solution was diluted in order to obta in an optica l dens ity be low 0.1 to avoid selfabsorption phenomena. The pulsed laser source is OPO from EKSPLA model NT230-50-SH, the excitation wavele ngth is 405nm w ith a repetition rate of 50Hz and a pulse width of 3ns. The emitted light was collected by an optical fiber placed at 90 degree angle from the excitation beam and fed to a monochromator from Jobin-Yvon (TRIAX 320) coupled to an EMCCD (Newton 920U from Andor). The sample is not placed at focus point. The laser spot diameter is measured as 2mm, and the power used here is from 0.1mW to 105mW.

A2.1 Spherical quantum wells solution sample

In FigureA2. 1, we present the norma lized PL spectra under varying power exc itation for 15nm, 25nm and 40nm SQWs. With the power increases, we observed the emergence of blue shoulder for 15nm SQWs , which is cons istent w ith the signature of multiexc itons in the time-resolved spectra, similarly, an emission feature can be seen at 2.4eV corresponding to the shell-related transition in the presence of a large number of excitons. In addition, different multie xcitons features are distinguishable on the blue shoulder in the larger particles compared to that in 15nm, which is the result of the larger absorption cross section, the emiss ion band at 2.0eV and 2.15eV is cons istence with the multicarriers(>4) involved S-level and P-level transition respectively.. However, when we plot the evolution of integrated PL intens ity as a function of excitation fluence, as shown in Figure A2. 1(b, d, f), the slope is around 0.5, the increase of the intensity is sublinear with <N> for the power range we used. The minimum of <N> here is ~15 exciton/pulse for 15nm SQWs, therefore plenty of multiexc itons formed in the SQWs, the sublinear indicated the existence of the non-radiative recombination, which is not completely suppressed. When the power continue to increase, the PL intensity continues to increase, suggesting that part of multiexcitons still recombined radiatively. This is the result of the high quantum yield of multiexcitons. For 40 nm SQWs, at 20 mJ/cm 2 , the curves start to saturate, this can be interpreted with the ultra large absorption cross-section, leading to the formation of much higher order excitons, in which the Auger recombination become predominant.

A2.2 Heterostructural nanoplatelets solution sample

We performed the same measurement on NPLs as shown in Figure A2. Annexe3: The influence of the applied voltage on PL spectra in nanoplatelets

In order to investigate the influence of electric field on photoluminescence, we applied an external electric field to CdSe/CdZnS NPLs according to ref. [163]. We focus on the change of photoluminescence intens ity and s pectral profile by adjusting the applied voltage. The schematic dia gram of the device is show n in Figure A3. 1. We designed a sandwich-like structure, it cons ists of two electrodes, which are made of FTO coated glass, and a NPLs flim was sandwished between the electrodes. Specifically, 100μl NPLs in polymer solution (PMAO) was deposited on the FTO surface of glass by spin-coating (1500rpm for 40s). The second FTO e lectrode was attached to the NPLs conta ining polymeric film w ith epoxy glue, the n dry for one night before applying the voltage. A DC voltage power supply is used to generate the external voltage. The experimenta l setup is based on the steady-state spectrum measurement setup mentioned previously. The applied voltage we used here is from 0V to 200V. The change of the integrated PL intensity and the spectra profile are presented in the Figure A3.2. The emission was effected strongly as the voltage increases. First, a great decrease of the emission intensity with increasing applied voltage is observed, which can be attributed to the charging effect of the NPLs [164].

We also observe a redshift of the peak position accompanied by a significant linewidth broadening. The spectra l broadening as a function of applied voltage is re lated to the increased LO (longitudinal optical)-phonon coupling in the presence of an electric field, while the red shift in peak position has been demonstrated as a characteristic of the quantum c onfined Stark effect. In ensembles of NCs , the stark shift a pproximate ly exhibits a quadratic dependence to the applie d voltage [164], which is consiste nt with our fitting in FigureA3.3.

Figure 2 . 1

 21 Figure 2.1 Schematic representation of qua ntum confinement effect: a bulk semiconductor w ith continue energy band a nd a fixe d energy gap E g,bulk , while NCs possess discrete energy levels and the band gap energy increases with decreasing size. . 20

Figure 2 . 21 Figure 2 . 3 22 Figure 2 . 4 25 Figure 2 . 5 26 Figure 2 . 6 29 Figure 2 . 7 31 Figure 2 . 8 32 Figure 2 . 9 33 Figure 2 . 34 Figure 2 . 35 Figure 2 .

 22123222425252626292731283229332342352 Figure 2.2 (a) Theoretical calculation on size dependence of the electron and hole energy levels in a CdSe quantum dots(QDs). (b) The comparison between t heoretical and experimenta l data for several of the lowest trans itions. ................................................. 21 Figure 2.3 Schematic diagram of the density of states as a function of dimensionality. Bulk materials(3D) without quantum confinement effects in three dimensions; 2D refers to quantum wells with quantum confinement effects in only one dimension;1D refers to rod-shaped or wire-shaped nanocrysta ls restricted in tw o dimens ions; 0D ma inly refers to quantum dots confined in three dimens ions. ................................................................. 22 Figure 2.4 Schematic representation of the energy level alignment and charge carrier localization in type I, quasi-type II and type II core/shell hetero-structural NCs, respectively................................................................................................................. 25 Figure 2.5 Schematic representation of poss ible relaxation and recombination processes in NCs after photoexc itation. (1) exciton is generated by absorption of a photon with energy ħꞷ E g or more photons are absorbed to form a multiexc itons state. (2)The hot exciton may relax to band edge by cooling process. (3)If ħꞷ 2E g , two or more excitons can be created through carrier multiplication(or multiple exciton generation, MEG) . (4)Multiexcitons will recombine radiatively to emit photon or non-ra diative ly via Auger recombination. (5) Eventually, the exciton recombines, relaxing the NCs back to the ground state. ............................................................................................................... 26 Figure 2.6 Schemes show ing the difference between Auger recombination process in direct gap bulk (left) and NCs (right) semiconductors. ∆E and ∆p present the change of energy and momentum, respective ly. E c and E v are the edge of conduction and va lence bands. Compared to bulk materia l, the requirement of mome ntum conservation is large ly relaxed in NCs and Auger recombination occurs effic iently for band edge carriers. ....... 29 Figure 2.7 Sche matic diagram of the exc iton and exciton Coulomb interaction in nanocrystals: (a) attractive force, (b) replusive force. ................................................... 31 Figure 2.8 Schematic diagram of the different radiative recombination pathways for triexciton in CdSe quantum dots. (a) The triexciton can decay by S-S electron-hole recombination, leading to an excited biexciton, following by the thermal process, (b) or by P-P recombination, leading to a ground-state biexciton, (c) or via the nominally

Figure 2 .

 2 Figure 2.13 (a)Schematic illustration of spherical quantum wells; (b) Quasi type II energy band a lignment of spherical quantum we lls. ................................................................. 42

Figure 2 . 42 Figure 2 .

 2422 Figure 2.14 (a)Abs orption a nd photoluminescence spectra of different she ll thickness SQWs(r: CdS core; l: CdSe well; h: CdS shell), (b) The evolution of ensemble PL decay dynamics with shell thickness (inset: single-exciton radiative recombination lifetime; R is the total size). ............................................................................................................. 42 Figure 2.15 Absorption and photoluminescence spectra of CdSe colloidal NPLs with different vertical thicknesses. ...................................................................................... 44

Figure 2 . 46 Figure 3 . 1 48 Figure 3 . 2 49 Figure 3 . 3 50 Figure 3 . 4 52 Figure 3 . 5 53 Figure 3 . 6 54 Figure 3 . 7 55 Figure 3 . 8 56 Figure 3 . 9 . 58 Figure 3 . 61 Figure 3 . 62 Figure 3 . 63 Figure 3 . 63 Figure 3 . 64 Figure 3 . 65 Figure 3 . 67 Figure 3 .

 2463148324933503452355336543755385639583613623633633643653673 Figure 2.16 (a)Schematic illustration of core/crown, core/shell and core/crown/shell heterostructures. (b) Absorption and photoluminescence spectra of CdSe/CdS core/crown, CdSe/CdZnS core/shell and CdSe/CdS/CdZnS core/crown/shell NPLs, all of which are synthesized at our laboratories. .................................................................................... 46 Figure 3.1 Red curve is the total Gibbs free energy as a function of the nuclei radius. ∆G c is the homogeneous nucleation energy barrier, and r c is the c r it ic a l rad ius. The bu lk termΔG 1 that scales with r 3 is plotted in green. The interfacial energy ΔG 2 is t he b lue curve that is proportional to surface area...................................................................... 48 Figure 3.2 Schematic illustration of the nucleation and subsequent growth process. ...... 49 Figure 3.3 Schematic of three neck flask apparatus for nanocrysta ls synthes is. .............. 50 Figure 3.4 Evolution of the absorption spectra with increasing s ize of CdS/CdSe quantum dots , which are synthes ized in our laboratory. .............................................................. 52 Figure 3.5 Evolution of the photoluminescence spectrum with increasing size of CdS/CdSe quantum dots, whic h are synthes ized in our laboratory. ................................ 53 Figure 3.6 Comparis on between ABS and PLE spectra. (a) without secondary nucleation and (b)w ith secondary nuc leation. (c) taken at different emiss ion wavelengths (the

Figure 3 . 70 Figure 3 . 71 Figure 3 . 72 Figure 3 . 73 Figure 3 . 74 Figure 3 . 75 Figure 3 .

 3703713723733743753 Figure 3.18 (a)Absorption spectra of CdSe/CdS core/crown NPLs with the crown increasing, (b) photoluminescence spectra of C/C NPLs synthesized by three methods with different crown s ize , comparing with core 4ML NPLs. ......................................... 70 Figure 3.19 TEM images of core/crown NPLs synthesized by three different approaches, (a) anisotropic growth solution, (b) S-ODE, (c) TOPS. (d) PLE of core/crown NPLs(approach 3) at different emiss ion wave lengths (as shown in inset). ..................... 71 Figure 3.20 Abs orption and photoluminescence spectra evolution of three core/shell CdSe/CdZnS NPLs with the increasing shell thickness. ................................................ 72 Figure 3.21 Photoluminescence excitation spectra of core/shell CdSe/CdZnS NPLs with different thickness shell measured at different emiss ion wave lengths. ........................... 73 Figure 3.22 TEM images of different thickness core/shell CdSe/CdZnS NPLs. (a) and (b) are core/shell_s, the arrow shows the vertical thickness is around 3.5nm, (c) and (d) are core/shell_m and core/shell_l, respectively................................................................... 74

Figure 4 . 1 80 Figure 4 . 2 81 Figure 4 . 3 82 Figure 4 . 4 84 Figure 4 . 5

 418042814382448445 Figure 4.1 Schematic representation of radiolumine scence and photoluminescence setup. .................................................................................................................................. 80Figure 4.2 Schematic representation of photoluminescence decay curve measurement setup........................................................................................................................... 81 Figure 4.3 Sche matic representation of time resolved emiss ion s pectrum setup under Xrays excitation ............................................................................................................ 82 Figure 4.4 Schematic representation of the multiexc itons formation process in nanocrysta ls (a) Absorbing three low energy photons simultaneous ly by inte nse pulsed excitation, (b) Simplifie d energy relaxation process by X-rays excitation (only three electron and hole pairs are shown). .............................................................................. 84Figure 4.5 Schematic energy ba nd dia gram and dens ity of state function of the SQWs. . 84

Figure 4 . 85 Figure 4 . 7 86 Figure 4 . 8 87 Figure 4 . 9 88 Figure 4 .

 4854786488749884 Figure 4.6 (a)Energy loss function restricted to 0-25 eV for CdS and CdSe. (b)Distribution of the number of scatterings for a 30keV primary electron for various partic le s izes. .............................................................................................................. 85 Figure 4.7 Simulation of evolution of e xcitations in NCs. (a) calc ulated time evolution of the population of neutral excitations (1, 2, 3 ... electron-hole pairs) for a 25nm SQWs considering an initial situation with a single secondary electron hole pair. Dashed curves corresponds to the simulation w ithout cons idering ra diative recombination. (b) and (c) Same simu la t io n w ith in it ia l s it uat io n w ith 2 , 3 secon dary e le ctron ho le pa irs. F or the simulation, radiative recombination considers a radiative decay time: τ X =50ns , τ BX =12. 5ns , τ TX =5.6ns. Excitation source : 20keV photons. ........................................... 86 Figure 4.8 Photoluminescence spectra comparison of solution samples (black line) and film samples (blue line). .............................................................................................. 87 Figure 4.9 Steady-state PL and RL spectra of a series of drop-cast SQWs (15nm, 25nm, 30nm, 40nm). Excitation source: 405nm light emitting diode(black line) and X-rays illumination at a voltage of 35 kV (red line). ................................................................ 88 Figure 4.10 Emiss ion intens ity evolution w ith different X-rays illumination time of 15nm and 40nm SQWs. The X-ray images were recorded at a voltage of 35 kV. The accumulation time is 400s. .......................................................................................... 89

Figure 4 .

 4 Figure 4.11 (a) PL decays of ensemble CdS/CdSe/CdS SQWs with varying shell thicknesses under low intens ity excitation w ithout focus ing and wave length filter. (b) Single exciton recombination lifetime for varying sized SQWs. .................................... 90

Figure 4 .

 4 Figure 4.12 Decay dynamics measured at three different energy position of a drop-cast 15nm dia meter SQWs. (a) <N> ~ 0.003, (b) <N> ~ 4. Black curve in (b) is the instrument response function (IRF).............................................................................. 91

Figure 4 . 93 Figure 4 . 94 Figure 4 .

 4934944 Figure 4.14 (a) Multie xcitons dynamics in 15nm SQWs obtained by subtracting the scaled single exciton contribution from high power decay (<N>~4), it can be fitted with two expone ntia ls by 0.72ns and 3.14ns. (b) Power dependent integra l PL inte ns ity of fast component(red square) and s low component(purple square). Slow component shows a single exciton s ignature and the fast component follows an approximately quadratic dependence on power. ................................................................................................. 93 Figure 4.15 The PL decay curves of 25nm(a) and 40nm(b) diameter SQWs under various excitation powers are norma lized at 100ns and 250ns to match the s ingle exc iton ta il, compared to that of 15nm SQWs, the multiexcitons components are longer lived. ......... 94 Figure 4.16 Comparison of time gated spectra corresponding to different time interva ls of 15nm SQWs. (a) 9.3nW, (b) 35.78μW.......................................................................... 96

Figure 4 .

 4 Figure 4.17 (a) Normalized time-resolved PL spectra in first 800ps as a function of power for drop-cast 15nm SQWs. The excitation wave length is 405nm. (b) Absorption (red), and absorption second derivative (black) spectra of 15nm SQWs. The arrows mark the 1S e -1S h 1.96eV and 1P e -1P h 2.16eV transitions. ..................................................... 97

Figure 4 . 98 Figure 4 . 99 Figure 4 . 100 Figure 4 .Figure 4 .Figure 4 .Figure 5 . 1 108 Figure5. 2 109 Figure 5 . 3 110 Figure 5 . 4 111 Figure 5 . 5 Figure 5 . 6 113 Figure 5 . 7 114 Figure 5 . 8

 4984994100444511082109531105411155561135711458 Figure 4.18 (a, c) Comparis on of spectra corresponding to different time interva ls of 25nm and 40nm SQWs at power 35.78μW. (b) The energy shift of different peak with respect to single exciton peak for 15nm, 25nm and 40nm SQWs. (d) Delayed emission at the late time of 40nm SQWs. ....................................................................................... 98 Figure 4.19 Integrated tota l photons under low(black line) and high(red line) power excitation of 15nm, 25nm and 40nm SQWs. ................................................................ 99 Figure 4.20 Decay dynamics measured at different energy of 15nm sample. Excitation source: X-rays illumination at a voltage of 30 kV. ...................................................... 100 Figure 4.21 Comparison of time gated spectra corresponding to different time interva ls of 15nm, 25nm and 40nm SQWs under X-rays exc itation. Exc itation source: X-rays illumination at a voltage of 30 kV. ............................................................................. 101 Figure 4.22 Comparison of time gated spectra at 0-2ns interva l for 15nm, 25nm and 40nm SQWs under X-rays exc itation w ith high intense photoexc itation. Exc itation source: X-rays illumination at a voltage of 30 kV and the laser power is 35.78μW....... 102 Figure 4.23 Comparison of the decay dynamics under X-rays and photoe xcitation of 15nm, 25nm a nd 40nm SQWs(defocus ing). Red curves are the multie xponentia l fitting. ................................................................................................................................ 103 Figure 5.1 Comparison of the optical properties between core/shell NPLs film(black line) and core/shell NPLs in polymer film(red line). (a) PL spectra recorded under low intensity optical excitation. (b) PL decay curves of the peak wavelength excited by low intens ity pulse d laser to avoid the formation of multiexc itons. .................................... 108

Figure 5 .

 5 Figure 5.9 (a-c) Power dependent PL decay trace of core/shell NPLs with varying shell thickness, the traces are normalized at long tail to match the single exciton, highlighting the fast components arising at intense optical excitation. (d) The fitting lifetimes of the fast component as a function of power. .......................................................................117

Figure 5 .Figure 5 .Figure 5 . 123 Figure 5 .Figure 5 .Figure 5 . 126 Figure 5 .

 5551235551265 Figure 5.10 (a)Decay dynamics at main peak of core/shell_ l NPLs for varying power intens ity. (b) The evolution of PL intens ity as a function of power: s low component is extracted from (a) at 60ns and fast component is the PL inte ns ity at 0ns after subtracting the contribution of single exc iton. ...............................................................................119 Figure 5.11 RL and PL decay curves of core/shell NPLs with varying shell thickness. (a,c,e) norma lized at long time , highlighting the multie xcitons compone nt. (b,d,f) The fast components of different sized core/shell NPLs for X-rays exc itation, obta ined by subtracting the single-exciton decay. The red curve in (b,d,f) shows biexponential fitting. ................................................................................................................................ 120 Figure 5.12 Normalized time gated RL(a,c,e) and PL (b,d,f) of different time intervals of core/shell CdSe/CdZnS NPLs with varying she ll thic kness. (a, b) core/shell_s, (c ,d) core/shell_m, (e,f) core/shell_l. X-rays illumination at a voltage of 30 kV voltage and 0.3μA current. The laser intensity is ~33μW............................................................... 123

138 FigureA2. 2 FigureA2. 3 FigureA3. 1 142 FigureA3. 3

 1382311423 Figure 5.17(a) Decay dyna mics at main peak of C/C/S NPLs for varying power inte ns ity. (b) The evolution of PL intens ity as a function of power: slow component is extracted from (a) at 75ns and fast component is the PL inte ns ity at 0ns after subtracting the contribution of single exc iton. ................................................................................... 127 Figure 5.18 (a) RL and PL decay curves of core/crown/she ll CdSe/CdS/CdZnS NPLs well norma lized at long time , highlighting the multiexc itons componen t. (b) Biexc iton fitting of fast compone nt under X-rays exc itation, obtained by subtracting the decay curve under photoexc itation from (a). ........................................................................ 128 FigureA2.1 (a,c,e)Emiss ion spectra of SQWs with varying she ll thic kness under varying laser power. (b,d,f) The evolution of integrated PL intens ity as a function of exc itation power. ...................................................................................................................... 138 FigureA2.2 Emiss ion s pectra of CdSe/CdS nanoplate lets under varying laser power.(b) The evolution of integrated PL intens ity as a function of exc itation power................... 139 FigureA2.3 Emiss ion spectra of CdSe/CdZnS nanoplatelets w ith varying she ll thickness under varying laser power. ........................................................................................ 139 FigureA3.1 Schematic representation of the device for applying voltage on NPLs: CdSe/CdZnS NPLs in polymer is sandw ishe d between two FTO electrodes coated on glass, epoxy glue is used to attach the second e lectrode. ............................................. 141 FigureA3.2 (a) integrated PL spectra under applie d volta ge from 0V to 200V, (b) Normalized voltage-dependent PL spectra, to highlight to linewidth broadening and the red shift of the peak position as the increasing applied voltage.................................... 142 FigureA3.3 The red shift of PL peak pos ition over applied voltage with respect to PL peak at 0V extracted from the spectra shown in Figure A3.2(b), fitted approximately with a quadratic function(~1.8) of applie d voltage (red line) .............................................. 142

Figure 4 . 117 Table 5 . 4 118 Table 5 . 5

 41175411855 23. ............. 104 Table 5.1. Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/crown NPLs under laser and X-rays exc itation. The average lifetimes: τ =(ΣA × τ )/ΣA , were calculated by using the respective amplitudes of each component....................................................................................113 Table 5. 2 Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/she ll NPLs with different shell thic kness under low intens ity photoexc itation and the calculated average lifetimes......................................116 Table 5. 3 The fitting parameters for tri-exponential decay for core/she ll_s NPLs under varying power excitation. ...........................................................................................The fitting parameters for tri-exponential decay for core/shell_m NPLs under varying power excitation. ...........................................................................................The fitting parameters for tri-exponential decay for core/she ll_ l NPLs under varying power excitation. ...........................................................................................118 Table 5. 6 Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/shell NPLs with different shell thickness under X-rays excitation and the calculated average lifetimes. .......................................................... 121 Table 5.7 Summary of the fitting parameters for multi-exponential decay, lifetimes with corresponding amplitudes for core/crown/shell NPLs under low inte ns ity photoexc itation and X-rays excitation, and the calculated average lifetimes......................................... 128

Figure 2 . 1

 21 Figure 2.1 Schematic representation of quantum confinement effect: a bulk semiconductor with continue energy band and a fixed energy gap Eg,bulk, while NCs possess discrete energy levels and the band gap energy increases with decreasing size.

  Figure 2.2(a) shows the electron and hole energy level structure of CdSe NCs with different sizes. 1S e , 1P e , 1D e and 2S e are the electron energy levels, and 1S 3/2 , 1P 3/2 and 2S 3/2 are the hole energy levels. It can be seen from the figure that the band gap of NCs increases as the size decreases. Since the effective mass of CdSe NCs holes is much larger than the effective mass of electrons, the electron energy level spacing is larger, and the hole energy leve ls are relative ly tight. More over, for several lowest transitions , the experimenta l results of energy changes w ith partic le size are highly cons iste nt with theoretical calculations, as shown in Figure 2(b).

Figure 2 .

 2 Figure 2.2 (a) Theoretical calculation on size dependence of the electron and hole energy levels in a CdSe quantum dot (QDs). (b) The comparison between theoretical and experimental data for several of the lowest transitions[20].

Figure 2 . 3

 23 Figure 2.3 Schematic diagram of the density of states as a function of dimensionality. Bulk materials(3D) without quantum confinement effects in three dimensions; 2D refers to quantum wells with quantum confinement effects in only one dimension;1D refers to rod-shaped or wireshaped nanocrystals restricted in two dimensions; 0D mainly refers to quantum dots confined in three dimensions.

Figure 2 . 4

 24 Figure 2.4 Schematic representation of the energy level alignment and charge carrier localization in type I, quasi-type II and type II core/shell hetero-structural NCs, respectively[30].

Figure 2 . 5

 25 Figure 2.5 Schematic representation of possible relaxation and recombination processes in NCs after photoexcitation. (1) exciton is generated by absorption of a photon with energy ħꞷ ≥ Eg or more photons are absorbed to form a multiexcitons state. (2)The hot exciton may relax to band edge by cooling process. (3)If ħꞷ ≥ 2Eg, two or more excitons can be created through carrier multiplic ation(or multiple exc iton generation, MEG) . (4)Multiexc itons w ill rec ombine radiatively to emit photon or non-radiatively via Auger recombination. (5) Eventually, the exciton recombines, relaxing the NCs back to the ground state[31].

Figure 2 . 6

 26 Figure 2.6 Schemes showing the difference between Auger recombination process in direct gap bulk (left) and NCs (right) semiconductors. ∆ and ∆ present the change of energy and momentum, respectively. Ec and Ev are the edge of conduction and valence bands. Compared to bulk material, the requirement of momentum conservation is largely relaxed in NCs and Auger recombination occurs efficiently for band edge carriers[18].

Figure 2 . 7

 27 Figure 2.7 Schematic diagram of the exciton and exciton Coulomb interaction in nanocrystals: (a) attractive force, (b) replusive force[48].

Figure 2 . 8

 28 Figure 2.8 Schematic diagram of the different radiative recombination pathways for triexciton in CdSe quantum dots. (a) The triexciton can decay by S-S electron-hole recombination, leading to an excited biexciton, following by the thermal process, (b) or by P-P recombination, leading to aground-state biexciton, (c) or via the nominally forbidden 1Pe-1Sh, also followed by a thermal process. Finally, each of these three pathways can bring the system to the ground-state biexciton.

Figure 2 . 9

 29 Figure 2.9 Schematic diagram of the carrier multiplication process. (a)Absorbing a photon with energy<2Eg, the electron dissipate the energy as heat. (b) After the absorption of a single highenergy photon(>2Eg), two excitons can be created through the Coulomb interaction in carrier multiplic ation proc es s s[39].

Figure 2 .

 2 Figure 2.10 Schematic representation of X-rays tube

Figure 2 .

 2 Figure 2.11 A s implified s c hematic repres entation of s cintillation detec tor[60] .

Figure 2 .

 2 Figure 2. 12 Sc hematic repres entation of s cintillation mec hanis m in bulk inorganic s c intillators . High energy photons absorbed can generate hot charge carriers, followed by thermalization to form a great amount of excitons and subsequent transport to traps involved intermediate states or luminescence centers, where radiative recombination occurs[64].

Figure 2 .

 2 Figure 2.13 (a)Sc hematic illus tration of s pheric al quantum w ells ; (b) Quas i type II energy band alignment of spherical quantum wells.

  Figure 2.14(b)).

Figure 2 .

 2 Figure 2.14 (a)Absorption and photoluminescence spectra of different shell thickness SQWs(r: CdS core; l: CdSe well; h: CdS shell), (b) The evolution of ensemble PL decay dynamics with s hell thic knes s (ins et: s ingle-exc iton radiative rec ombination lifetime; R is the total s ize)[81].

Figure 2 .

 2 Figure 2.15 Absorption and photoluminescence spectra of CdSe colloidal NPLs with different vertical thicknesses[91].

Figure 2 .

 2 Figure 2.16 (a)Sc hematic illus tration of c ore/c rown, c ore/shell and c ore/c rown/s hell heterostructures. (b) Absorption and photoluminescence spectra of CdSe/CdS core/crown, CdSe/CdZnS core/shell and CdSe/CdS/CdZnS core/crown/shell NPLs, all of which are synthesized at our laboratories.

Figure 3 . 1

 31 Figure 3.1 Red curve is the total Gibbs free energy as a function of the nuclei radius. ∆Gc is the homogeneous nucleation energy barrier, and rc is the critical radius. The bulk termΔG1 that scales with r 3 is plotted in green. The interfacial energy ΔG2 is the blue curve that is proportional to surface area.

Figure 3 .

 3 Figure 3.2 schematically illustrated the processes of nucleation and subsequent growth. Typica lly nanopartic les are synthesized in three stages[102]. The concentration of monomer increases with time in stage I, no nucleation would occur even above the equilibrium solubility because of the extremely high energy barrier for homogeneous nucleation. The nucleation stage II appears only when the super-saturation reaches the critica l nuc leation threshold, thereby overcoming the energy barrier, which corresponds to the free energy defined by equation 3-4. While as the formation of nuc le i, the concentration of monomer decreases until the critical energy, nucleation stage ends and the reaction enter into growth stage until the concentration of monomer reaches the equilibr iu m solubility. Lamer theoretic al mechanism dia gram is shown below.

Figure 3 . 2

 32 Figure 3.2 Sc hematic illus tration of the nuc leation and s ubs equent grow th proces s.

Figure 3 . 3

 33 Figure 3.3 Schematic of three neck flask apparatus for nanocrystals synthesis.

Figure 3 . 4

 34 Figure 3.4 Evolution of the absorption spectra with increasing size of CdS/CdSe quantum dots, which are synthesized in our laboratory. Size is calculated by the precursors used

Figure 3 . 5

 35 Figure 3.5 Evolution of the photoluminescence spectrum with increasing size of CdS/CdSe quantum dots, which are synthesized in our laboratory. Size is calculated by the precursors used.

Figure 3 . 6

 36 Figure 3.6 Comparison between ABS and PLE spectra. (a) without secondary nucleation and (b)with secondary nucleation. (c) taken at different emission wavelengths (the maximum of the emission, lower energy side and the higher energy side of the emission peak).

Figure 3 . 7

 37 Figure 3.7 Schematic representation of absolute method with integrating sphere. Left: substrate spectrum measurement, right: sample spectrum measurement.

Figure 3 . 8

 38 Figure 3.8 TEM images of nanoparticles at different magnifications: A.20000X, B.120000X, C.300000X.

1

 1 Chemical materials -1-octadecene, 90%,CAS : 112-88-9. Sigma Aldrich -1-dodecanethiol, 98%, CAS : 112-55-0, Sigma Aldrich -cadmium nitrate tetrahydrate, >99%, CAS : 10022-68-1, Sigma-Aldrich -cadmium oxide, 99.99%, CAS : 1306-19-0, Strem -ethanol,>=99%, CAS : 64-17-5, Carlo Erba -hexane, CAS : 110-54-3, Carlo Erba -isopropanol, 99.9%, CAS : 67-63-0, VWR -methanol, 98.5%, CAS : 67-56-1, VWR -oleic acid, 90%, CAS : 112-80-1, Alfa-aesar -oleylamine, 80-90%, CAS : 112-90-3, Acros -selenium, 99.99%, CAS : 7782-49-2, Strem -sulfur, 99.99%, CAS : 7704-34-9, Strem -sodium myristate, >98%, CAS : 822-12-8, TCI -toluene, 99.9%, CAS : 108-88-3, VWR -trioctylphosphine , 97%, CAS : 4731-53-7, Sigma-Aldrich

Figure 3 . 9 .

 39 Figure 3.9. The protocol of spherical quantum wells synthesis

Figure 3 .

 3 Figure 3.10 Absorption spectra of CdS seeds in hexane

Figure 3 .

 3 Figure 3.11 (a)Absorption spectrum of CdS/CdSe in hexane: variety color curves represent different aliquots taken during size growing, showing a red shift; (b)TEM image of CdS/CdSe at 400000X magnification, inset is the histogram of size distribution of ~100 particles.

Figure 3 .

 3 Figure 3.12 Absorption(left) and photoluminescence(right) spectra of CdS/CdSe and different sized spherical quantum wells: 15nm, 20nm, 25nm 30nm, 35nm, 40nm.

Figure 3 .

 3 Figure 3.13 Image of different size SQWs from 15nm to 40nm dispersed in hexane or toluene with the same concentration under white light and UV light excitation.

Figure 3 .

 3 Figure 3.14 Comparison between ABS and PLE spectra. (a) CdS/CdSe and the same sample with additional ligands measured at maximum emission; (b) spherical quantum wells (30nm in diameter).

Figure 3 .

 3 Figure 3.15 Transmission electron microscopy (TEM) images of different size SQWs at magnification 30000X, insets are histograms of size distribution.

Figure 3 .

 3 Figure 3.16 Absorption and normalized photoluminescence spectra of 4ML nanoplatelets. The excitation photon wavelength is 350nm.

Figure 3 .

 3 Figure 3.17 TEM images of 4ML CdSe NPLs.

Figure 3 .

 3 Figure 3.18 (a)Absorption spectra of CdSe/CdS core/crown NPLs with the crown increasing, (b) photoluminescence spectra of C/C NPLs synthesized by three methods with different crown size, comparing with core 4ML NPLs.

Figure 3 .

 3 Figure 3.19 TEM images of core/crown NPLs synthesized by three different approaches, (a) anisotropic growth solution, (b) S-ODE, (c) TOPS. (d) PLE of core/crown NPLs(approach 3) at different emission wavelengths (as shown in inset).

Figure 3 .

 3 Figure 3.20 Absorption and photoluminescence spectra evolution of three core/shell CdSe/CdZnS NPLs w ith the inc reas ing s hell thic knes s .

Figure 3 .

 3 Figure 3.21 Photoluminescence excitation spectra of core/shell CdSe/CdZnS NPLs with different thickness shell measured at different emission wavelengths.

Figure 3 .

 3 22(b) and (c).

Figure 3 .

 3 Figure 3.22 TEM images of different thicknesses core/shell CdSe/CdZnS NPLs. (a) and (b) are core/shell_s, the arrow shows the vertical thickness is around 3.5nm, (c) and (d) are core/shell_m and core/shell_l, respectively.

3. 4 . 5

 45 Hot-injection core/crown/shell CdSe/CdS/CdZnS

Figure 3 .

 3 Figure 3.23 (a)Absorption and photoluminescence spectra comparison of core/shell (black dash line) and core/crown/shell NPLs (red solid line). (b)PLE of C/C/S NPLs with different thickness shell measured at different emission wavelengths.

Figure 3 .

 3 Figure 3.24 TEM image of core/crown/shell CdSe/CdS/CdZnS NPLs.

Figure 4 . 1

 41 Figure 4.1 Schematic representation of radioluminescence and photoluminescence setup.

Figure 4 . 2

 42 Figure 4.2 Schematic representation of photoluminescence decay curve measurement setup.

Figure 4 . 3

 43 Figure 4.3 Schematic representation of time resolved emission spectrum setup under X-rays excitation

Figure 4 . 4

 44 Figure 4.4 Schematic representation of the multiexcitons formation process in nanocrystals (a) Absorbing three low energy photons simultaneously by intense pulsed excitation, (b) Simplified energy relaxation process by X-rays excitation (only three electron and hole pairs are shown).

4. 2 . 2

 22 Simula tio n of the scintilla tion process in nanocrystals This s imulation part has been fully performed by Prof. Andrey Vas il'ev from the Lomonosov Moscow State University. In order to simulate the physical mechanisms occurring a long the energy relaxation, we used the micro-theory of scintillation fully based on the die lectric response function of the materia l, which is recently observed in the reference[139]. Origina lly deve loped for s ingle crysta ls, the mode l has been adapted for nanocrysta ls , namely cons idering the geometrical parameters of the SQWs and introducing a work function in order to build up the energy diagram (Figure 4.5).

Figure 4 . 5

 45 Figure 4.5 Schematic energy band diagram and density of state function of the SQWs.

Figure 4 .

 4 Figure 4.6 (a)Energy loss function restricted to 0-25 eV for CdS and CdSe. (b)Distribution of the number of scatterings for a 30keV primary electron for various particle sizes.

Figure 4 .

 4 6(a) (about 12 eV for CdS a nd CdSe). The mean number of SEs produced by an electron pass ing through a NC is therefore (-/ )∆ /ℏ, where (-/ ) is the stopping power (2.5 eV/nm for an electron having a kinetic energy of 30keV in CdS), ∆ the mean free path of an electron through NC with diameter d ( ∆ =2d/3) from geometrical considerations). The number of SEs created in matter should be distributed according to Poisson distribution. More accurate consideration demands techniques like the one used by Landau when he described so-called Landau fluctuations, i.e. fluctuation of energy depos ited by a partic le in thin layer with ∆ width. Just for the estimation, the average of the SEs with energy ℏ produced by a fast electron in NC with diameter d is distributed in nearly Poisson way as seen in F igure 4.6(b). As an intermediate remark, the distribution of the number of SEs highlights that multi-SE is likely to occur even in the 15nm partic les and is dominating at larger sizes. For that reason, we have considered for the second steps of the simulation several initia l situations where the partic les conta in 1, 2, 3... SE with energies distributed along the energy loss function.

Figure 4 . 7

 47 Figure 4.7 Simulation of evolution of excitations in NCs. (a) calculated time evolution of the population of neutral excitations (1, 2, 3 ... electron-hole pairs) for a 25nm SQWs considering an initial situation with a single secondary electron hole pair. Dashed curves corresponds to the simulation without considering radiative recombination. (b) and (c) Same simulation with initial situation with 2, 3 secondary electron hole pairs. For the simulation, radiative recombination considers a radiative decay time: τX=50ns, τBX=12.5ns, τTX=5.6ns. Excitation source: 20keV photons.

Figure 4 . 8

 48 Figure 4.8 Photoluminescence spectra comparison of solution samples (black line) and film s amples (blue line).

Figure 4 . 9

 49 Figure 4.9 Steady-state PL and RL spectra of a series of drop-cast SQWs (15nm, 25nm, 30nm, 40nm). Excitation source: 405nm light emitting diode(black line) and X-rays illumination at a voltage of 35 kV (red line).

Figure 4 .

 4 Figure 4.10 Emis s ion intens ity evolution w ith different X-rays illumination time of 15nm and 40nm SQWs. The X-ray images were recorded at a voltage of 35 kV. The accumulation time is 400s.

Figure 4 .

 4 Figure 4.11 (a) PL decays of ensemble CdS/CdSe/CdS SQWs with varying shell thicknesses under low intensity excitation without focusing and wavelength filter. (b) Single exciton recombination lifetime for varying sized SQWs.

Figure 4 .

 4 Figure 4.12 Decay dynamics measured at three different energy positions of a drop-cast 15nm diameter SQWs. (a) <N> ~ 0.003, (b) <N> ~ 4. Black curve in (b) is the instrument response function (IRF).

Figure4. 13

 13 Figure4.[START_REF] Kamkaew | Scintillat ing Nanoparticles as Energy Mediators for Enhanced Photodynamic Therapy[END_REF] The PL decay curves of 15nm diameter SQWs under various excitation powers are normalized at 60ns to match the single exciton tail, highlighting the fast multiexcitons component at early time.

Figure 4 .

 4 Figure 4.14 (a) Multiexcitons dynamics in 15nm SQWs obtained by subtracting the scaled single exciton contribution from high power decay (<N>~4), it can be fitted with two exponentials by 0.72ns and 3.14ns. (b) Power dependent integral PL intensity of fast component (red square) and slow component (purple square). Slow component shows a single exciton signature and the fast component follows an approximately quadratic dependence on power.

Figure 4 .

 4 Figure 4.15 The PL decay curves of 25nm(a) and 40nm(b) diameter SQWs under various excitation powers are normalized at 100ns and 250ns to match the single exciton tail, compared to that of 15nm SQWs, the multiexcitons components are longer lived.

Figure 4 .

 4 Figure 4.16 Comparison of time gated spectra corresponding to different time intervals of 15nm SQWs. (a) 9.3nW, (b) 35.78μW.

Figure 4 . 17

 417 Figure 4.17(a) Normalized time-resolved PL spectra in first 800ps as a function of power for drop-cast 15nm SQWs. The excitation wavelength is 405nm. (b) Absorption (black), and absorption second derivative (red) spectra of 15nm SQWs. The arrows mark the 1Se-1Sh 1.96eV and 1Pe-1Ph 2.16eV transitions.

  Figure 4.17(a) Normalized time-resolved PL spectra in first 800ps as a function of power for drop-cast 15nm SQWs. The excitation wavelength is 405nm. (b) Absorption (black), and absorption second derivative (red) spectra of 15nm SQWs. The arrows mark the 1Se-1Sh 1.96eV and 1Pe-1Ph 2.16eV transitions.

Figure 4 .

 4 Figure 4.18 (a, c) Comparison of spectra corresponding to different time intervals of 25 nm and 40 nm SQWs at power 35.78 μW. (b) The energy shift of the different peaks with respect to single exciton peak for 15 nm, 25 nm and 40 nm SQWs. (d) Delayed emission at the late time of 40 nm SQWs.

Figure 4 .

 4 Figure 4.19 presents the comparison of total emitted photons between low(black line)and high(red line) power excitation, which are consistent with the steady state PL. As we expected, the spectral profile under high power excitation is analogue to that under Xrays excitation, that is , a shoulder is observed on the blue edge , involving the formation of multiexcitons and, as the shell thickness increases, the spectrum under intense photoexc itation shows a slight blue shift around 3 meV, 6 meV and 12 meV for 15 nm, 25 nm and 40 nm SQWs, respectively. These observations reveal the energy recombination under X-rays radiation might undergo the similar process with that of inte nse photoexc itation. Meanwhile , it a lso points out that to compare the photoexc itation is a feasible way to investigate the excited state in radioluminescence. Therefore, in the following section we will further give the decay curve and time-resolved RL spectrum under X-rays excitation, and then explore the composition of its emitted light through comparative analysis.

Figure 4 .

 4 Figure 4.19 Integrated total photons under low(black line) and high(red line) power excitation of 15nm, 25nm and 40nm SQWs.

Figure 4 .

 4 Figure 4.20 Decay dynamics measured at different energy of 15nm sample. Excitation source: Xrays illumination at a voltage of 30 kV.

Figure 4 .

 4 Figure 4.21 Comparison of time gated spectra corresponding to different time intervals of 15nm, 25nm and 40nm SQWs under X-rays excitation. Excitation source: X-rays illumination at a voltage of 30 kV.

Figure 4 .

 4 Figure 4.22 Comparison of time gated spectra at 0-2ns interval for 15nm, 25nm and 40nm SQWs under X-rays excitation with high intense photoexcitation. Excitation source: X-rays illumination at a voltage of 30 kV and the laser power is 35.78μW.

Figure 4 .

 4 Figure 4.23 Comparison of the decay dynamics under X-rays and photoexcitation of 15nm, 25nm and 40nm SQWs(defocusing). Red curves are the multiexponential fitting.

Figure 5 . 1

 51 Figure 5.1 Comparison of the optical properties between core/shell NPLs film(black line) and core/shell NPLs in polymer film(red line). (a) PL spectra recorded under low intensity optical excitation. (b) PL decay curves of the peak wavelength excited by low intensity pulsed laser to avoid the formation of multiexcitons.

Figure5. 2

 2 Figure5.2 Radiolumines c enc e s pec tra of C/S NPLs for different illuminated time. Exc itation source: X-rays generated by 35kV voltage and 35mA current.

5. 2

 2 Core/crown CdSe/CdS nanoplatelets 5.2.1 Steady state radioluminescence

Figure 5 . 3

 53 Figure 5.3 The photoluminescence spectrum is fitted by two different functions. (a) Lorentzian fitting, (b) Gaussian fitting.

Figure 5 . 4

 54 Figure 5.4 The comparison between PL and RL spectra and their fitting line shape with pseudo-Voigt function. PL was measured at low power to avoid the multiexcitons generation. (a) PL spectrum shows a single symmetric peak. And the RL spectrum is fitted by the sum of two components, y=A1V1+A2V2, the corresponding fitting of V1(blue line) and V2(green line) are also shown in the Figure(b).

Figure 5 . 5

 55 Figure 5.5 Normalized time resolved RL spectra of different time intervals, early time (0-0.8 ns)and later time (4-6 ns) spectra reveals a red-shifted spectrum at early times, pointing toward multiexcitons generation under X-rays excitation. X-rays illumination at a voltage of 30 kV voltage and 0.3μA current.

Figure 5 . 6

 56 Figure 5.6 Normalized time-resolved fluorescence decay curves of CdSe/CdS core/crown NPLs under photoexcitation (23μW/cm 2 ) (black circle) and X-rays excitation(blue circle), red line is the tri-exponential fitting curve. Photoexcitaion is measured by defocusing of 405nm pulsed laser(repetition:500kHz) to avoid the multiexcitons.

Figure 5 . 7

 57 Figure 5.7 Normalized steady state PL and RL spectra comparison of core/shell CdSe/CdZnS NPLs with varying shell thickness.

Figure 5 . 8

 58 Figure 5.8 Normalized fluorescence decay curves of CdSe/CdZnS core/shell NPLs with different shell thickness in polymer under low intensity photoexcitation(defocusing), red lines are the triexponential fitting curves.

Figure 5 .

 5 Figure 5.9 (a-c) Power dependent PL decay trace of core/shell NPLs with varying shell thickness, the traces are normalized at long tail to match the single exciton, highlighting the fast components arising at intense optical excitation. (d) The fitting lifetimes of the fast component as a function of power.

Figure 5 .

 5 Figure 5.10 (a)Decay dynamics at main peak of core/shell_l NPLs for varying power intensity. (b)The evolution of PL intensity as a function of power: slow component is extracted from (a) at 60ns and fast component is the PL intensity at 0ns after subtracting the contribution of single exciton.

Figure 5 .

 5 Figure 5.11 RL and PL decay curves of core/shell NPLs with varying shell thickness. (a,c,e) normalized at long time, highlighting the multiexcitons component. (b,d,f) The fast components of different sized core/shell NPLs for X-rays excitation, obtained by subtracting the singleexciton decay. The red curve in (b,d,f) shows biexponential fitting.

Figure 5 .

 5 Figure 5.12 Normalized time gated RL(a,c,e) and PL (b,d,f) of different time intervals of core/shell CdSe/CdZnS NPLs with varying shell thickness. (a, b) core/shell_s, (c,d) core/shell_m, (e,f) core/shell_l. X-rays illumination at a voltage of 30 kV voltage and 0.3μA current. The laser intens ity is ~33μW.

Figure 5 .

 5 Figure 5.13 Normalized steady state PL and RL spectra for core/crown/shell CdSe/CdS/CdZnS NPLs. The power of photoexciton source is low with 405nm wavelength.

Figure 5 .

 5 Figure 5.14 Normalized PL spectra with different power intensity for C/C/S NPLs solution sample. Excitation source: OPO with 405nm wavelength and 50Hz.

Figure 5 .

 5 Figure 5.15 Normalized time gated PL for C/C/S NPLs in different time intervals as indicated in the figures under different power intensity. (a) Laser power is 4nW, (b) and (d) the slices for 0-60ns and after 60ns at laser power 33μW. (c) The position of the main peak and the change of the spectral width over time at laser power 33μW.

Figure 5 .

 5 Figure 5.16 Normalized RL spectra for C/C/S NPLs in different time intervals as indicated in the figures under X-rays excitation with 30kV accelerating voltage.

Figure 5 . 17

 517 Figure 5.17(a) Decay dynamics at main peak of C/C/S NPLs for varying power intensity. (b) The evolution of PL intensity as a function of power: slow component is extracted from (a) at 75ns and fast component is the PL intensity at 0ns after subtracting the contribution of single exciton.

  Figure 5.17(a) Decay dynamics at main peak of C/C/S NPLs for varying power intensity. (b) The evolution of PL intensity as a function of power: slow component is extracted from (a) at 75ns and fast component is the PL intensity at 0ns after subtracting the contribution of single exciton.

Figure 5 .

 5 18 (b). Through biexponential fitting, we can extract the lifetime of the fast component about 0.5ns. The specific fitting parameters are shown in Table5.7.

Figure 5 .

 5 Figure 5.18 (a) RL and PL decay curves of core/crown/shell CdSe/CdS/CdZnS NPLs well normalized at long time, highlighting the multiexcitons component. (b) Biexciton fitting of fast component under X-rays excitation, obtained by subtracting the decay curve under photoexcitation from (a).

FigureA2. 1

 1 FigureA2.1 (a,c,e)Emission spectra of SQWs with varying shell thickness under varying laser power. (b,d,f) The evolution of integrated PL intensity as a function of excitation power.

  2 and A2.3, we summarized the behavior of different heterostructural NPLs under OPO excitation. A2.2.1 Core/crown CdSe/CdS

FigureA2. 2

 2 FigureA2.2 Emission spectra of CdSe/CdS nanoplatelets under varying laser power.(b) The evolution of integrated PL intensity as a function of excitation power.

FigureA2. 3

 3 FigureA2.3 Emission spectra of CdSe/CdZnS nanoplatelets with varying shell thickness under varying laser power.

FigureA3. 1

 1 FigureA3.1 Schematic representation of the device for applying voltage on NPLs: CdSe/CdZnS NPLs in polymer is sandwished between two FTO electrodes coated on glass, epoxy glue is used to attach the second electrode.

FigureA3. 2

 2 FigureA3.2 (a) integrated PL spectra under applied voltage from 0V to 200V, (b) Normalized voltage-dependent PL spectra, to highlight to linewidth broadening and the red shift of the peak position as the increasing applied voltage.

FigureA3. 3

 3 FigureA3.3 The red shift of PL peak position over applied voltage with respect to PL peak at 0V extracted from the spectra shown in Figure A3.2(b), fitted approximately with a quadratic function(~1.8) of applied voltage (red line)

Table 3 .

 3 1 The reaction parameters for SQWs shell growth. ........................................... 60

Table 4 .

 4 [START_REF] Murray | Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites[END_REF] The ratio of various emiss ion species w ith the neutra l X. ............................... 87

Table 4 .

 4 2 Exponentia l fitting results and exc itonic spec ies assignment of 15nm, 25nm a nd 40nm SQWs. .............................................................................................................. 94

Table 4 .

 4 3 Summary of the fitting parameters for decay dynamics under X-rays excitation of 15nm, 25nm and 40nm SQWs. .............................................................................. 104

Table 4 .

 4 [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF] The ratio of different exc itation species to s ingle exc iton for different s ized SQWs, the data extracted from the fitting of RL decay curves in

Table A1 .

 A1 1 Absorption cross-section estimations for different sized SQWs ................. 135 Table A1.2 The average number of exc iton forme d per nanocrysta l per pulse(<N>) for different sized SQWs ................................................................................................ 136

	Chapter 1
	List of abbreviations ABS Absorption ASE Amplified spontaneous emission BX Biexciton Cd(Ac) 2 Cadmium acetate dihydrate Cd(myr) 2 Cadmium myristate CdO Cadmium oxide Cd(OA) 2 Cadmium oleate CM Carrier multiplication DDT 1-dodecanethiol FRET Forster resonance energy transfer FWHM Full-width-at-half-maximum MEG Multiple exciton generation MX Multiexciton NPLs Nanoplatelets OAm Oleylamine OA Oleic acid ODE 1-octadecene PL Photoluminescence PLE Photoluminescence excitation PMAO Poly maleic anhydride-alt-octadecene QY Quantum yield QD Quantum dot RL Radioluminescence SE Secondary excitation SILAR Successive ion layer adsorption and reaction S-ODE Sulfur-in-octadecene solution SQWs Spherical quantum wells TEM Transmission electron microscope TOP Trioctylphospine TOPS Sulfur-in-trioctylphos pine solution TOPSe Selenium-in-trioctylphospine solution TX Triexciton Zn(Ac) 2 Zinc acetate dihydrate Zhu MENG Chapter 1

Table 3 .

 3 [START_REF] Murray | Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites[END_REF] The reaction parameters for SQWs shell growth.

	Diame te r(nm)	15	20	25*	30	35	40
	Cd precursor (mL)	19	33	43	51	63	72
	S pre cursor (mL)	19	33	43	51	63	72
	Ratio of aliquot	1/2	1	1/3	1/3	1/3	1
	Concentration( ) # 3.5	2.6	0.87	0.57	0.38	0.77

*: the reaction was ended at 20nm SQWs, and sample was washed and dispersed in 10mL hexane. The continuous reaction was restarted based on 20nm SQWs(1/2 of total volume ) in 10mL ODE. #: concentration is calculated for SQWs dispersed in 10mL hexane.

Table 4 .

 4 [START_REF] Murray | Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites[END_REF] The ratio of various emission species with the neutral X.

Table 4 .

 4

2 Exponential fitting results and excitonic species assignment of 15 nm, 25 nm and 40 nm SQWs.

Table 4 .

 4 3 Summary of the fitting parameters for decay dynamics under X-rays excitation of 15nm, 25nm and 40nm SQWs.

		(ns)	(ns)	(ns)	(ns)
	15nm	29	8.5	2.6	0.6
	25nm	75	18.3	3.6	0.6
	40nm	180	35	4.8	0.7

Table 4 .

 4 [START_REF] Lim | Optical gain in colloidal quantum dots achieved with direct-current electrical pumping[END_REF] The ratio of different excitation species to single exciton for different sized SQWs, the data extracted from the fitting of RL decay curves in Figure4.23.

		15nm	25nm	40nm
	X * /X	0.29	0.50	0.55
	BX/X	0.22	0.30	0.33
	MX/X	0.13	0.13	0.11
	Table

Table 5 .

 5 2 Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/shell NPLs with different shell thickness under low intensity photoexcitation and the calculated average lifetimes.

		1(ns )	A1	2(ns )	A2	3(ns )	A3	avg(ns )
	core/shell_s	2.57	0.56	10.56	0.41	52.88	0.04	7.63
	core/shell_m	3.44	0.53	13.4	0.44	60.01	0.04	10.34
	core/shell_l	12.24	0.38	33.74	0.48	103.05	0.1	32.55

Table 5 .

 5

	Power(μW)	1(ns )	A1	2(ns )	A2	3(ns )	A3	4(ns )	A4
	0.07	33.2	0.02	7.89	0.14	1.61	0.12
	0.101	33	0.02	7.16	0.16	1.19	0.16
	0.177	33	0.01	9.42	0.11	3.11	0.15	0.72	0.2
	0.333	33	0.01	8.82	0.12	2.46	0.17	0.57	0.25
	0.8	33	0.02	8.12	0.14	1.96	0.21	0.43	0.37

[START_REF] García-Santamaría | Suppressed Auger Recombination in "Giant" Nanocrystals Boosts Optical Gain Performance[END_REF]

The fitting parameters for tri-exponential decay for core/shell_s NPLs under varying power excitation.

Table 5 . 4

 54 The fitting parameters for tri-exponential decay for core/shell_m NPLs under varying power excitation.

	Power(μW)	1(ns )	A1	2(ns )	A2	3(ns )	A3	4(ns )	A4
	0.07	39.87	0.02	11. 24	0.16	2.63	0.13
	0.101	40	0.02	12.56	0.14	3.93	0.15	0.89	0.1
	0.177	40	0.02	11. 14	0.16	2.77	0.21	0.56	0.18
	0.333	40	0.02	11. 43	0.15	3.06	0.2	0.68	0.28
	0.8	40	0.03	10.08	0.19	2.2	0.29	0.44	0.46

Table 5 . 5

 55 The fitting parameters for tri-exponential decay for core/shell_l NPLs under varying power excitation.

	Power(μW)	1(ns )	A1	2(ns )	A2	3(ns )	A3	4(ns )	A4
	0.07	94.83	0.02	22.83	0.16	4.49	0.12
	0.101	94	0.02	25.44	0.13	7.88	0.12	1.77	0.15
	0.177	94	0.01	25.58	0.14	6.31	0.18	1.23	0.3
	0.333	94	0.02	23.96	0.16	4.94	0.22	0.87	0.36
	0.8	94	0.02	20.86	0.19	3.21	0.3	0.52	0.48

Table 5 .

 5 [START_REF] Röntgen | On a New Kind of Rays[END_REF] Summary of the fitting parameters for tri-exponential decay, lifetimes with corresponding amplitudes for core/shell NPLs with different shell thickness under X-rays excitation and the calculated average lifetimes.

		1(ns ) A1	2(ns ) A2	3(ns ) A3	4(ns )	A4	avg(ns )
	core/shell_s	0.51	6.28	2.84	1.74 12.95	0.4	52	<0.01	1.59
	core/shell_m 0.54	8.3	2.87	2.16	14.4	0.45	60	<0.01	1.57
	core/shell_l	0.64	9.81	3.56	2.71	23.1	0.83	100	0.13	3.58

1(ns ) A1 2(ns ) A2 3(ns ) A3 4(ns ) A4 avg(ns )

	photoexcitat ion			5.36	0.36 29.84 0.46	160	0.12	55.2
	X-rays	0.5	12.01	3.46	3.23 25.61	0.9	160	0.12	3.65

Table A1 .

 A1 1 Absorption cross-section estimations for different sized SQWs

	Diame te r (nm)	15	20	25	30	35	40
	(cm 2 )	8× 10 -14	1.9× 10 -13	3× 10 -13	6.4× 10 -13	1× 10 -12	1.2× 10 -12

Table A1 .

 A1 [START_REF] Schaller | High Efficiency Carrier Multiplication in PbSe Nanocrystals: Implications for Solar Energy Conversion[END_REF] The average number of exciton formed per nanocrystal per pulse(<N>) for different sized SQWs

	Power(nW) 1	9.3	29.5	77	264	702 1654 4399 12739 35789
	15nm	0.003 0.024 0.074 0.19 0.67 1.78 4.18 11. 1 32.2	90.5
	25nm	0.009 0.09	0.28	0.71 2.51 6.68 15.7 41.6 120	339
	40nm	0.04	0.36	1.11	2.85 10.1 26.7 62.7 166	483	1357

Scientific backgroundZhu MENG

Scientific background