# Comprehensive investigation of Er$_{2}$O$_{3}$ thin films grown with different ALD approaches

L. Khomenkova, H. Merabet, M.-P. Chauvat, C. Frilay, X. Portier, C. Labbe, P. Marie, Julien Cardin, S. Boudin, J.-M. Rueff, et al.

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L. Khomenkova, H. Merabet, M.-P. Chauvat, C. Frilay, X. Portier, et al. Comprehensive investigation of Er$_{2}$O$_{3}$ thin films grown with different ALD approaches. Surfaces and Interfaces, 2022, 34, pp.102377. 10.1016/j.surfin.2022.102377. hal-03790146

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# Comprehensive investigation of Er₂O₃ thin films grown with different ALD approaches

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L. Khomenkova<sup>a,d,e,*</sup>, H. Merabet<sup>b</sup>, M.-P. Chauvat<sup>a</sup>, C. Frilay<sup>a</sup>, X. Portier<sup>a</sup>, C. Labbe<sup>a</sup>, P. Marie<sup>a</sup>, J. Cardin<sup>a</sup>, S. Boudin<sup>c</sup>, J.-M. Rueff<sup>c</sup>, F. Gourbilleau<sup>a,*</sup>

## ARTICLE INFO

Keywords:

Er₂O₃ thin films

Atomic layer deposition

TEM

XRD

Luminescence

Er silicate

## ABSTRACT

The effect of Er precursor nature (Er(CpMe)₃ or Er(tmhd)₃) and annealing treatment at 500–1100 °C on the structural and optical properties of Er₂O₃ films grown on Si substrates by thermal or O₂-plasma-assisted atomic layer deposition was studied by means of spectroscopic ellipsometry, Fourier-transform infrared spectroscopy, X-ray diffraction, transmission electron microscopy coupled with energy dispersive X-ray spectroscopy as well as photoluminescence method. An annealing at 500–800 °C resulted in the film crystallization mainly. Thermal treatment at high temperatures caused the formation of Er silicate phase due to the diffusion of Si atoms from the substrate in the films depth. This phase was found to be Er₂SiO₅ being crystallized at 1100 °C. Light emitting properties of the films are determined by Er₂O₃ native defects (like oxygen vacancies) and intra-4f shell transition in Er³⁺ ions. The latter dominated in the films annealed at 1000–1100 °C. The most intense Er³⁺ emission, observed in the films grown with O₂-plasma-assisted approach, was explained by a lower contribution of oxygen vacancies as well as by pronounced crystallization of Er silicate phase. In this latter, the effect of concentration quenching of Er³⁺ luminescence was lower due to a larger distance between Er³⁺ neighbor ions.

## 1. Introduction

Over the last decades, atomic layer deposition (ALD) is widely used due to continuous downscaling of electronic devices as well as to the development of 3D-integrating circuits [1,2]. This technique allows forming thin films of different materials not only on flat substrates, but also on complex 3D-surfaces and structures at relatively low temperatures [3,4]. The films with specific chemical composition can be easily deposited via alternating exposure of gaseous chemical precursors and reactants. Due to self-limiting saturated reactions, the thickness of the films and their composition can be controlled via repetition of different ALD cycles. The slow growth of ALD films, that is usually caused by an incomplete removing of large molecular groups or generated by-products, is no longer an hindrance as traditional coating processes because of fine tuning of film thickness at a sub-nanometer scale.

Nowadays, rare earth sesquioxides and materials on their basis

attract considerable attention because of their wide range of applications such as optical devices, microelectronics, magnetic devices, photonics and photovoltaics. These oxides are known to exist in three main polymorphs, i.e. hexagonal, monoclinic and cubic bixbyite structures (so-called, A-, B- and C-forms, respectively) [5]. Different research groups reported on the growth of these materials on different semiconducting substrates including C-Sc₂O₃/GaN [6,7], La₂O₃/Si [8], Pr₂O₃/Si [9], Sm₂O₃/Si [10], C-Er₂O₃/Si [11,12], C-Er₂O₃/GaN [13], Yb₂O₃/Si [14], etc.

It is known that Er₂O₃ has good heat endurance and chemical durability under high temperatures that makes it attractive for the protection and anticorrosion coating of steels [15]. Due to its high dielectric constant, Er₂O₃ is applicable also for gate insulators [16,17]. Owing to intense photoluminescence at room temperature, Er₂O₃ materials can be used for the fabrication of high-gain optical waveguide amplifiers and optoelectronic integrated circuits [18,19].

<footnote><sup>a</sup> CIMAP, CNRS, ENSICAEN, UNICAEN, CEA, Normandie Université, 6 Boulevard Maréchal Juin, Caen 14050 CEDEX 4, France</footnote>

<footnote><sup>b</sup> Physics Program, Department of Mathematics, Statistics, and Physics, College of Arts and Sciences, Qatar University, P.O. Box 2713, Doha, Qatar</footnote>

<footnote><sup>c</sup> Normandie Univ., ENSICAEN, UNICAEN, CNRS, CRISMAT, 14000 Caen, France</footnote>

<footnote><sup>d</sup> V. Lashkaryov Institute of Semiconductor Physics at the National Academy of Sciences of Ukraine, 45 Pr.Nauky, Kyiv 03028, Ukraine</footnote>

<footnote><sup>e</sup> National University of Kyiv-Mohyla Academy, 2 Skovorody str., Kyiv 04070, Ukraine</footnote>

<footnote>Abbreviations: TEM, Transmission electron microscopy; XRD, X-ray diffraction; ALD, Atomic layer deposition.</footnote>

<footnote>* Corresponding authors at: CIMAP, ENSICAEN, UNICAEN, CEA, CNRS, Normandie Université, 6 Boulevard Maréchal Juin, Caen 14050 CEDEX 4, France.
E-mail addresses: khomen@ukr.net (L. Khomenkova), fabrice.gourbilleau@ensicaen.fr (F. Gourbilleau).</footnote>
# Comprehensive investigation of Er₂O₃ thin films grown with different ALD approaches

<figure><img src="image_2.png" /></figure>

L. Khomenkova<sup>a,d,e,*</sup>, H. Merabet<sup>b</sup>, M.-P. Chauvat<sup>a</sup>, C. Frilay<sup>a</sup>, X. Portier<sup>a</sup>, C. Labbe<sup>a</sup>, P. Marie<sup>a</sup>, J. Cardin<sup>a</sup>, S. Boudin<sup>c</sup>, J.-M. Rueff<sup>c</sup>, F. Gourbilleau<sup>a,*</sup>

## ARTICLE INFO

**Keywords:**

Er₂O₃ thin films

Atomic layer deposition

TEM

XRD

Luminescence

Er silicate

## ABSTRACT

The effect of Er precursor nature (Er(CpMe)₃ or Er(tmhd)₃) and annealing treatment at 500–1100 °C on the structural and optical properties of Er₂O₃ films grown on Si substrates by thermal or O₂-plasma-assisted atomic layer deposition was studied by means of spectroscopic ellipsometry, Fourier-transform infrared spectroscopy, X-ray diffraction, transmission electron microscopy coupled with energy dispersive X-ray spectroscopy as well as photoluminescence method. An annealing at 500–800 °C resulted in the film crystallization mainly. Thermal treatment at high temperatures caused the formation of Er silicate phase due to the diffusion of Si atoms from the substrate in the films depth. This phase was found to be Er₂SiO₅ being crystallized at 1100 °C. Light emitting properties of the films are determined by Er₂O₃ native defects (like oxygen vacancies) and intra-4f shell transition in Er³⁺ ions. The latter dominated in the films annealed at 1000–1100 °C. The most intense Er³⁺ emission, observed in the films grown with O₂-plasma-assisted approach, was explained by a lower contribution of oxygen vacancies as well as by pronounced crystallization of Er silicate phase. In this latter, the effect of concentration quenching of Er³⁺ luminescence was lower due to a larger distance between Er³⁺ neighbor ions.

## 1. Introduction

Over the last decades, atomic layer deposition (ALD) is widely used due to continuous downscaling of electronic devices as well as to the development of 3D-integrating circuits [1,2]. This technique allows forming thin films of different materials not only on flat substrates, but also on complex 3D-surfaces and structures at relatively low temperatures [3,4]. The films with specific chemical composition can be easily deposited via alternating exposure of gaseous chemical precursors and reactants. Due to self-limiting saturated reactions, the thickness of the films and their composition can be controlled via repetition of different ALD cycles. The slow growth of ALD films, that is usually caused by an incomplete removing of large molecular groups or generated by-products, is no longer an hindrance as traditional coating processes because of fine tuning of film thickness at a sub-nanometer scale.

Nowadays, rare earth sesquioxides and materials on their basis

attract considerable attention because of their wide range of applications such as optical devices, microelectronics, magnetic devices, photonics and photovoltaics. These oxides are known to exist in three main polymorphs, i.e. hexagonal, monoclinic and cubic bixbyite structures (so-called, A-, B- and C-forms, respectively) [5]. Different research groups reported on the growth of these materials on different semiconducting substrates including C-Sc₂O₃/GaN [6,7], La₂O₃/Si [8], Pr₂O₃/Si [9], Sm₂O₃/Si [10], C-Er₂O₃/Si [11,12], C-Er₂O₃/GaN [13], Yb₂O₃/Si [14], etc.

It is known that Er₂O₃ has good heat endurance and chemical durability under high temperatures that makes it attractive for the protection and anticorrosion coating of steels [15]. Due to its high dielectric constant, Er₂O₃ is applicable also for gate insulators [16,17]. Owing to intense photoluminescence at room temperature, Er₂O₃ materials can be used for the fabrication of high-gain optical waveguide amplifiers and optoelectronic integrated circuits [18,19].

<footnote><sup>a</sup> CIMAP, CNRS, ENSICAEN, UNICAEN, CEA, Normandie Université, 6 Boulevard Maréchal Juin, Caen 14050 CEDEX 4, France</footnote>

<footnote><sup>b</sup> Physics Program, Department of Mathematics, Statistics, and Physics, College of Arts and Sciences, Qatar University, P.O. box 2713, Doha, Qatar</footnote>

<footnote><sup>c</sup> Normandie Univ., ENSICAEN, UNICAEN, CNRS, CRISMAT, 14000 Caen, France</footnote>

<footnote><sup>d</sup> V. Lashkaryov Institute of Semiconductor Physics at the National Academy of Sciences of Ukraine, 45 Pr.Nauky, Kyiv 03028, Ukraine</footnote>

<footnote><sup>e</sup> National University of Kyiv-Mohyla Academy, 2 Skovorody str., Kyiv 04070, Ukraine</footnote>

<footnote>Abbreviations: TEM, Transmission electron microscopy; XRD, X-ray diffraction; ALD, Atomic layer deposition.</footnote>

<footnote>* Corresponding authors at: CIMAP, ENSICAEN, UNICAEN, CEA, CNRS, Normandie Université, 6 Boulevard Maréchal Juin, Caen 14050 CEDEX 4, France.
E-mail addresses: khomen@ukr.net (L. Khomenkova), fabrice.gourbilleau@ensicaen.fr (F. Gourbilleau).</footnote>
The methods of Er₂O₃ thin film deposition include physical vapor deposition [18,20,21], chemical vapor deposition [11,22], sol-gel method [23], magnetron sputtering [11,24,25] and, more recently, ALD technique [12,26].

In this work, Er₂O₃ films were grown on Si substrates by either thermal or plasma-assisted ALD approaches. Hereafter, the effect of deposition conditions and post-deposition treatments on the optical, structural and luminescent properties of these films will be compared and discussed.

## 2. Experimental details

The Er₂O₃ thin films were grown using two Er precursors, i.e. Er (CpMe)₃ (tris(methylcyclopentadienyl)erbium(III) (99.9%-Er), CAS No. 39470-10-5, chemical formula Er(CH₃C₅H₄)₃, formula weight 404.62 g/mol) and Er(tmhd)₃ (tris(2,2,6,6-tetramethyl-3,5-heptanedionato) erbium(III), 99% (99.9%-Er), CAS No. 35733-23-4, chemical formula Er(C₁₁H₁₉O₂)₃, formula weight 717.08 g/mol). Both precursors were provided by STREM Chemicals Inc. [27].

### 2.1. Properties of Er precursors

Prior to install Er precursors in the ALD deposition tool, they were submitted to thermogravimetric analysis (TGA) using a SETERAM TAG 92 setup. The measurements were done under air atmosphere at a heating rate of 3°C/min from room temperature to 800°C. The TG curves obtained for both Er precursors, Er(CpMe)₃ and Er(tmhd)₃, at atmospheric pressure show different profiles (Fig. 1,a). The Er(CpMe)₃ precursor exhibits a continuous weight loss in the 25–800 °C temperature range with a final weight loss of 51.7 % which is in good agreement with the theoretical value of 52.7% expected for the formation of Er₂O₃.

Regarding the Er(tmhd)₃ precursor, this material shows a three-step degradation process with temperature. Between 25 and 170 °C the TG curve presents a plateau showing that the precursor is stable in this temperature range. Above 170 °C and up to 318 °C, a first weight loss of 87% is observed immediately followed by a second mass loss of 7.25% up to 800 °C. Contrary to the Er(CpMe)₃ compound, the final weight loss

of 94.25 % for Er(tmhd)₃ at 800 °C is much more important than the theoretical value of 73.3% expected for the formation of Er₂O₃. This mass difference (Δm = 20.95%) can be explained by the concomitant decomposition and sublimation-vaporization of Er(tmhd)₃ which both contribute to the weight loss during TGA study due to the volatility of (tmhd)-chelates of rare earths [28].

The XRD analysis performed for both precursors before and after TG experiments showed that initial Er(tmhd)₃ precursor presents a crystalline structure whereas Er(CpMe)₃ is amorphous (Fig. 1,b). The final material obtained after TGA for both precursors (Fig. 1,b) was identified as Er₂O₃ oxide in its cubic phase [29,30].

It should be emphasized that thermal behavior of both Er precursors was extracted from the TG experiment performed in air atmosphere at ambient pressure. This fact can explain some difference in mass loss in comparison with the results reported for the similar precursors by other groups ([12,31]). The multi-step mass loss upon the heating is the evidence of precursors' decomposition via step-like desorption of corresponding organic ligands. Similar effect was previously described, for instance, for Cp-based Zr precursors [32,33] and explained by the difference in interactions in the gas phase in TG experiment and on the surface of condensed phase during the ALD process.

Successful deposition of the films using ALD technique depends on many parameters such as substrate temperature, pulse/purge duration of the precursor and the oxidant, flow of vector-gas, chamber pressure, etc. Their effect on the film growth is considered below.

### 2.2. Samples

Er₂O₃ thin films were grown using a Picosun® R-200 Advanced deposition tool from Picosun Oy company (Finland) [34]. P-type (B-doped) Si wafers grown by Czochralski pulling technique produced by "Sil'tronix Silicon Technologies" company (France) were used [35]. The wafers, being of (100) orientation and double side polished, had the following dimensions: 2 inches in diameter and 275 ± 25 µm in thickness. Their resistivity was ρ = 15 Ω·cm. The virgin substrates were covered with a 5-nm thermal SiO₂ layer.

Er(CpMe)₃ and Er(tmhd)₃ precursors were introduced in the ALD

<figure><img src="image_2.png" /><figcaption>**Fig. 1.** (a) Weight loss as a function of temperature for Er(tmhd)₃ and Er(CpMe)₃ precursors; (b) XRD patterns recorded on Er(tmhd)₃ (1,3) and Er(CpMe)₃ (2,4) precursors before (1,2) and after (3,4) thermogravimetric experiments.</figcaption></figure>
The methods of Er₂O₃ thin film deposition include physical vapor deposition [18,20,21], chemical vapor deposition [11,22], sol-gel method [23], magnetron sputtering [11,24,25] and, more recently, ALD technique [12,26].

In this work, Er₂O₃ films were grown on Si substrates by either thermal or plasma-assisted ALD approaches. Hereafter, the effect of deposition conditions and post-deposition treatments on the optical, structural and luminescent properties of these films will be compared and discussed.

## 2. Experimental details

The Er₂O₃ thin films were grown using two Er precursors, i.e. Er (CpMe)₃ (tris(methylcyclopentadienyl)erbium(III) (99.9%-Er), CAS No. 39470-10-5, chemical formula Er(CH₃C₅H₄)₃, formula weight 404.62 g/mol) and Er(tmhd)₃ (tris(2,2,6,6-tetramethyl-3,5-heptanedionato) erbium(III), 99% (99.9%-Er), CAS No. 35733-23-4, chemical formula Er(C₁₁H₁₉O₂)₃, formula weight 717.08 g/mol). Both precursors were provided by STREM Chemicals Inc. [27].

### 2.1. Properties of Er precursors

Prior to install Er precursors in the ALD deposition tool, they were submitted to thermogravimetric analysis (TGA) using a SETERAM TAG 92 setup. The measurements were done under air atmosphere at a heating rate of 3°C/min from room temperature to 800°C. The TG curves obtained for both Er precursors, Er(CpMe)₃ and Er(tmhd)₃, at atmospheric pressure show different profiles (Fig. 1,a). The Er(CpMe)₃ precursor exhibits a continuous weight loss in the 25–800 °C temperature range with a final weight loss of 51.7 % which is in good agreement with the theoretical value of 52.7% expected for the formation of Er₂O₃.

Regarding the Er(tmhd)₃ precursor, this material shows a three-step degradation process with temperature. Between 25 and 170 °C the TG curve presents a plateau showing that the precursor is stable in this temperature range. Above 170 °C and up to 318 °C, a first weight loss of 87% is observed immediately followed by a second mass loss of 7.25% up to 800 °C. Contrary to the Er(CpMe)₃ compound, the final weight loss

of 94.25 % for Er(tmhd)₃ at 800 °C is much more important than the theoretical value of 73.3% expected for the formation of Er₂O₃. This mass difference (Δm = 20.95%) can be explained by the concomitant decomposition and sublimation-vaporization of Er(tmhd)₃ which both contribute to the weight loss during TGA study due to the volatility of (tmhd)-chelates of rare earths [28].

The XRD analysis performed for both precursors before and after TG experiments showed that initial Er(tmhd)₃ precursor presents a crystalline structure whereas Er(CpMe)₃ is amorphous (Fig. 1,b). The final material obtained after TGA for both precursors (Fig. 1,b) was identified as Er₂O₃ oxide in its cubic phase [29,30].

It should be emphasized that thermal behavior of both Er precursors was extracted from the TG experiment performed in air atmosphere at ambient pressure. This fact can explain some difference in mass loss in comparison with the results reported for the similar precursors by other groups ([12,31]). The multi-step mass loss upon the heating is the evidence of precursors' decomposition via step-like desorption of corresponding organic ligands. Similar effect was previously described, for instance, for Cp-based Zr precursors [32,33] and explained by the difference in interactions in the gas phase in TG experiment and on the surface of condensed phase during the ALD process.

Successful deposition of the films using ALD technique depends on many parameters such as substrate temperature, pulse/purge duration of the precursor and the oxidant, flow of vector-gas, chamber pressure, etc. Their effect on the film growth is considered below.

### 2.2. Samples

Er₂O₃ thin films were grown using a Picosun® R-200 Advanced deposition tool from Picosun Oy company (Finland) [34]. P-type (B-doped) Si wafers grown by Czochralski pulling technique produced by “Sil’tronix Silicon Technologies” company (France) were used [35]. The wafers, being of (100) orientation and double side polished, had the following dimensions: 2 inches in diameter and 275 ± 25 µm in thickness. Their resistivity was ρ = 15 Ω·cm. The virgin substrates were covered with a 5-nm thermal SiO₂ layer.

Er(CpMe)₃ and Er(tmhd)₃ precursors were introduced in the ALD

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** (a) Weight loss as a function of temperature for Er(tmhd)₃ and Er(CpMe)₃ precursors; (b) XRD patterns recorded on Er(tmhd)₃ (1,3) and Er(CpMe)₃ (2,4) precursors before (1,2) and after (3,4) thermogravimetric experiments.</figcaption></figure>
reactor from outside containers with the flow of 150 sccm. Nitrogen (99.999 %) was used as a carrier and purge gas. The main deposition parameters are collected in Table 1S (see electronic supplement information).

Thermal ALD deposition was performed with Er(CpMe)₃ precursor and distilled water used as oxidant. The water container was kept at 25 °C. The temperature of the bottle with Er(CpMe)₃ was varied in the range T<sub>b</sub> = 120–200 °C. The value of T<sub>b</sub> = 200 °C was the highest reachable for bottle configuration in our deposition tool. The pulse/purge durations for Er(CpMe)₃ were varied in the range of t<sub>pulse</sub> = 1–3 s and t<sub>purge</sub> = 2–6 s. The precursor flow was 150 sccm. To increase the amount of the precursor in the reaction zone, the boost was used with the flow varied from 250 to 800 sccm. The water flow was kept as 150 sccm and pulse/purge durations were fixed at t<sub>pulse</sub> = 0.1 s and t<sub>purge</sub> = 4 s.

Plasma-assisted ALD process was carried out with Er(tmhd)₃ and oxygen plasma. A container with Er(tmhd)₃ was heated at T<sub>b</sub> = 155–200 °C. Corresponding Er(tmhd)₃ pulse/purge time were t<sub>pulse</sub> = 2–4 s and t<sub>purge</sub> = 4–6 s. Oxygen plasma was introduced in the reactor with a flow of 150 sccm and the pulse/purge times of t<sub>pulse</sub>=13–56 and t<sub>purge</sub> = 2–4 s. The radio-frequency plasma power (RFP) and its duration were varied in the 1500–3000 W and 10–52 s ranges, respectively.

The substrate temperature was tuned in the T<sub>d</sub> = 275–325 °C range for both approaches. The film thickness was controlled via the cycles' number.

After the evacuation of the substrates with uniformly deposited films from the ALD tool, they were cut into small pieces (called hereafter as samples) and submitted to annealing treatment at different temperatures (T<sub>A</sub> = 500–1100 °C) and durations (t<sub>A</sub> = 10–30 min) in nitrogen atmosphere.

## 2.3. Experimental techniques used for the film characterization

To investigate optical, structural and luminescent properties of the films versus annealing treatment, several techniques were used. The spectroscopic ellipsometry study was performed using a Jobin-Yvon ellipsometer (UVISEL, HORIBA Ltd., Kyoto, Japan). The incident light was scanned in the 1.5–6.0 eV range and incident angles were 65°, 70° and 75°. The beam spot was 1 mm in diameter. The spectra were recorded from the central point of each sample.

The surface of the films was analysed by means of optical microscopy using Olympus microscope and atomic-force microscopy (AFM) using the NanoScope® III Scanning probe microscope (Digital Instruments) operated in amplitude and phase tapping modes.

Vibrational properties of the films were investigated using Fourier transform infrared spectroscopy (FTIR). The FTIR absorption spectra were recorded with a Nicolet™ iS50 FTIR spectrometer with normal and 60° incidence of excitation light.

Structural properties of the films were studied using X-ray diffraction and X-ray diffraction methods. The XRD data were collected using X-ray diffractometer Philips X'Pert-MRD with the Cu Kα₁ radiation (λ = 0.15418 nm) or Co Kα₁ radiation (λ = 0.17889 nm). The patterns were recorded in Bragg-Brentano (θ-2θ) and grazing (ω-2θ) geometries.

For Transmission Electron Microscopy (TEM) observations, the cross sectional thin foils of the films were prepared by FIB (Focused Ion Beam) such that the TEM observations were done following a [110] direction of the (100) Si substrate. In such a way, the electron beam was parallel to the film/substrate interface. The FIB system used was a Dual beam FEI Helios nanolab660 with an electron imaging resolution of 0.6 nm at 15 kV (Field Emission Gun - FEG) and a FIB resolution of 2.54 nm at 30 kV. In order to protect the film surface from the gallium ion beam during the thinning process, a carbon deposit was carried out beforehand as well as two platinum layers deposited electronically (0.3 µm) and then ionically (2 µm). The thinning was processed by a set of repetitive inclinations of the gallium beam (±1° with decreasing gradually the beam currents in the first phases and a final surface cleaning step with a higher inclination

but for lower currents) on either side of the thin section until electron transparency (a few tens of nanometers). The conventional bright and dark fields (CTEM) and high resolution TEM (HRTEM) observations, scanning transmission electron microscopy and high angle annular dark field (STEM-HAADF) observations as well as STEM-EDX chemical imaging were performed using a JEOL ARM200F microscope with a “cold” field emission gun (FEG - 200 kV) and having the double corrections of aberrations in image and in STEM modes. The point to point resolutions were 0.1 nm in TEM mode and 0.078 nm in STEM-HAADF mode. This microscope was equipped with an EDX Centurio JEOL spectrometer. All digital images were processed with the GMS2 GATAN Digital Micrograph software.

Raman scattering spectra were recorded in a quasi-backscattering geometry using a Jobin Yvon LabRaman 300 spectrometer with a resolution of 1.57 cm⁻¹, equipped with a confocal microprobe and a CCD camera. As an excitation source, the second harmonic of Nd:YAG laser (532 nm) was used. The laser beam power was around 1.5 mW and the spot diameter was 0.87 µm. Under such conditions, no heating of the sample was observed during spectra recording. This setup was also used to collect micro-photoluminescence spectra. Besides, standard PL and PL Excitation (PLE) spectra were measured with Horiba Fluorolog-3 setup (model FL3-22). It is equipped with a 450-W Xenon light source, double-grating excitation and emission spectrometers of Czerny-Turner design with kinematic classically-ruled gratings and all-reflective optics, with automated slits, and room-temperature R928P detector. A 2-nm band-pass is achieved using 1200-grooves/mm gratings with a dispersion of 2.1 nm/mm and a 1 mm slits. Calibrated photodiode is used for excitation correction from 240-1000 nm. Emission detector is an R928P for high sensitivity in photon-counting mode in the 240–850 nm spectral range. To record Er³⁺ photoluminescence in infrared spectral range, a Jobin Yvon 1 m single grating monochromator equipped with Hamamtsu detector was used. The excitation was performed using a 488-nm line of Ar⁺ laser. The PL and PLE spectra were corrected automatically with dedicated software during spectra recording. The measurements were performed at room temperature.

# 3. Experimental results and discussion

## 3.1. Parameters of the films obtained from spectroscopic ellipsometry

Spectroscopic ellipsometry is known as a fast, sensitive and non-destructive method for the investigation of optical properties of thin films. Hereafter we used an approach described in details in our previous works [36–38]. In brief, a spectrum consists in the measured Ψ and Δ ellipsometric angles defined from the fundamental equation $\bar{r}_p/\bar{r}_s = \tan\Psi \cdot \exp(i\Delta)$, where $\bar{r}_p$ and $\bar{r}_s$ are the complex reflection coefficients for parallel and perpendicular polarizations of light, respectively. The light intensity in both polarizations is considered as $I_s = I \cdot \sin2\Psi \cdot \sin\Delta$ and $I_c = I \cdot \sin2\Psi \cdot \cos\Delta$. For ellipsometric modeling of a rough film, this latter is divided it into two parts: one part containing the homogeneous film and the other part representing a non-homogeneous form referred to as “roughness layer”. The first part is simulated with the physically meaningful dielectric function belonging to the material. To model the film roughness, Maxwell-Garnett effective medium approximation was applied. This approach allowed to obtain self-agreed results not only for as-deposited, but also for the films annealed, especially, at high temperatures. The roughness value of the films ($R_p$) estimated from one ellipsometric spectrum is considered as a mean height calculated over the entire measured surface illuminated by excitation light in one run that corresponding to the beam spot size.

Fig. 2,a shows typical spectra and their modeling applying a “new-amorphous” model described recently in Refs. [36–38]. The use of three different angles of incident light improves the accuracy for the determination of the film thickness, *d*, and film roughness, as well as refractive index, *n*, and the absorption coefficient, *α* (if any). By the
reactor from outside containers with the flow of 150 sccm. Nitrogen (99.999 %) was used as a carrier and purge gas. The main deposition parameters are collected in Table 1S (see electronic supplement information).

Thermal ALD deposition was performed with Er(CpMe)₃ precursor and distilled water used as oxidant. The water container was kept at 25 °C. The temperature of the bottle with Er(CpMe)₃ was varied in the range T<sub>b</sub> = 120–200 °C. The value of T<sub>b</sub> = 200 °C was the highest reachable for bottle configuration in our deposition tool. The pulse/purge durations for Er(CpMe)₃ were varied in the range of t<sub>pulse</sub> = 1–3 s and t<sub>purge</sub> = 2–6 s. The precursor flow was 150 sccm. To increase the amount of the precursor in the reaction zone, the boost was used with the flow varied from 250 to 800 sccm. The water flow was kept as 150 sccm and pulse/purge durations were fixed at t<sub>pulse</sub> = 0.1 s and t<sub>purge</sub> = 4 s.

Plasma-assisted ALD process was carried out with Er(tmhd)₃ and oxygen plasma. A container with Er(tmhd)₃ was heated at T<sub>b</sub> = 155–200 °C. Corresponding Er(tmhd)₃ pulse/purge time were t<sub>pulse</sub> = 2–4 s and t<sub>purge</sub> = 4–6 s. Oxygen plasma was introduced in the reactor with a flow of 150 sccm and the pulse/purge times of t<sub>pulse</sub>=13–56 and t<sub>purge</sub> = 2–4 s. The radio-frequency plasma power (RFP) and its duration were varied in the 1500–3000 W and 10–52 s ranges, respectively.

The substrate temperature was tuned in the T<sub>d</sub> = 275–325 °C range for both approaches. The film thickness was controlled via the cycles' number.

After the evacuation of the substrates with uniformly deposited films from the ALD tool, they were cut into small pieces (called hereafter as samples) and submitted to annealing treatment at different temperatures (T<sub>A</sub> = 500–1100 °C) and durations (t<sub>A</sub> = 10–30 min) in nitrogen atmosphere.

## 2.3. Experimental techniques used for the film characterization

To investigate optical, structural and luminescent properties of the films versus annealing treatment, several techniques were used. The spectroscopic ellipsometry study was performed using a Jobin-Yvon ellipsometer (UVISEL, HORIBA Ltd., Kyoto, Japan). The incident light was scanned in the 1.5–6.0 eV range and incident angles were 65°, 70° and 75°. The beam spot was 1 mm in diameter. The spectra were recorded from the central point of each sample.

The surface of the films was analysed by means of optical microscopy using Olympus microscope and atomic-force microscopy (AFM) using the NanoScope® III Scanning probe microscope (Digital Instruments) operated in amplitude and phase tapping modes.

Vibrational properties of the films were investigated using Fourier transform infrared spectroscopy (FTIR). The FTIR absorption spectra were recorded with a Nicolet™ iS50 FTIR spectrometer with normal and 60° incidence of excitation light.

Structural properties of the films were studied using X-ray diffraction and X-ray diffraction methods. The XRD data were collected using X-ray diffractometer Philips X'Pert-MRD with the Cu Kα₁ radiation (λ = 0.15418 nm) or Co Kα₁ radiation (λ = 0.17889 nm). The patterns were recorded in Bragg-Brentano (θ-2θ) and grazing (ω-2θ) geometries.

For Transmission Electron Microscopy (TEM) observations, the cross sectional thin foils of the films were prepared by FIB (Focused Ion Beam) such that the TEM observations were done following a [110] direction of the (100) Si substrate. In such a way, the electron beam was parallel to the film/substrate interface. The FIB system used was a Dual beam FEI Helios nanolab660 with an electron imaging resolution of 0.6 nm at 15 kV (Field Emission Gun - FEG) and a FIB resolution of 2.54 nm at 30 kV. In order to protect the film surface from the gallium ion beam during the thinning process, a carbon deposit was carried out beforehand as well as two platinum layers deposited electronically (0.3 µm) and then ionically (2 µm). The thinning was processed by a set of repetitive inclinations of the gallium beam (±1° with decreasing gradually the beam currents in the first phases and a final surface cleaning step with a higher inclination

but for lower currents) on either side of the thin section until electron transparency (a few tens of nanometers). The conventional bright and dark fields (CTEM) and high resolution TEM (HRTEM) observations, scanning transmission electron microscopy and high angle annular dark field (STEM-HAADF) observations as well as STEM-EDX chemical imaging were performed using a JEOL ARM200F microscope with a “cold” field emission gun (FEG - 200 kV) and having the double corrections of aberrations in image and in STEM modes. The point to point resolutions were 0.1 nm in TEM mode and 0.078 nm in STEM-HAADF mode. This microscope was equipped with an EDX Centurio JEOL spectrometer. All digital images were processed with the GMS2 GATAN Digital Micrograph software.

Raman scattering spectra were recorded in a quasi-backscattering geometry using a Jobin Yvon LabRaman 300 spectrometer with a resolution of 1.57 cm⁻¹, equipped with a confocal microprobe and a CCD camera. As an excitation source, the second harmonic of Nd:YAG laser (532 nm) was used. The laser beam power was around 1.5 mW and the spot diameter was 0.87 µm. Under such conditions, no heating of the sample was observed during spectra recording. This setup was also used to collect micro-photoluminescence spectra. Besides, standard PL and PL Excitation (PLE) spectra were measured with Horiba Fluorolog-3 setup (model FL3-22). It is equipped with a 450-W Xenon light source, double-grating excitation and emission spectrometers of Czerny-Turner design with kinematic classically-ruled gratings and all-reflective optics, with automated slits, and room-temperature R928P detector. A 2-nm band-pass is achieved using 1200-grooves/mm gratings with a dispersion of 2.1 nm/mm and a 1 mm slits. Calibrated photodiode is used for excitation correction from 240-1000 nm. Emission detector is an R928P for high sensitivity in photon-counting mode in the 240–850 nm spectral range. To record Er³⁺ photoluminescence in infrared spectral range, a Jobin Yvon 1 m single grating monochromator equipped with Hamamtsu detector was used. The excitation was performed using a 488-nm line of Ar⁺ laser. The PL and PLE spectra were corrected automatically with dedicated software during spectra recording. The measurements were performed at room temperature.

# 3. Experimental results and discussion

## 3.1. Parameters of the films obtained from spectroscopic ellipsometry

Spectroscopic ellipsometry is known as a fast, sensitive and non-destructive method for the investigation of optical properties of thin films. Hereafter we used an approach described in details in our previous works [36–38]. In brief, a spectrum consists in the measured Ψ and Δ ellipsometric angles defined from the fundamental equation $\bar{r}_p/\bar{r}_s = \tan\Psi \cdot \exp(i\Delta)$, where $\bar{r}_p$ and $\bar{r}_s$ are the complex reflection coefficients for parallel and perpendicular polarizations of light, respectively. The light intensity in both polarizations is considered as $I_s = I \cdot \sin2\Psi \cdot \sin\Delta$ and $I_c = I \cdot \sin2\Psi \cdot \cos\Delta$. For ellipsometric modeling of a rough film, this latter is divided it into two parts: one part containing the homogeneous film and the other part representing a non-homogeneous form referred to as “roughness layer”. The first part is simulated with the physically meaningful dielectric function belonging to the material. To model the film roughness, Maxwell-Garnett effective medium approximation was applied. This approach allowed to obtain self-agreed results not only for as-deposited, but also for the films annealed, especially, at high temperatures. The roughness value of the films ($R_p$) estimated from one ellipsometric spectrum is considered as a mean height calculated over the entire measured surface illuminated by excitation light in one run that corresponding to the beam spot size.

Fig. 2,a shows typical spectra and their modeling applying a “new-amorphous” model described recently in Refs. [36–38]. The use of three different angles of incident light improves the accuracy for the determination of the film thickness, *d*, and film roughness, as well as refractive index, *n*, and the absorption coefficient, *α* (if any). By the
<figure><img src="image_3.png" /><figcaption>**Fig. 2.** (a) Typical ellipsometry spectra for as-deposited Er₂O₃ film (symbols) taken at 65°, 70° and 75° and their fitting (solid curves) using "new amorphous" model. (b) Refractive index and absorption coefficient of as-deposited Er₂O₃ film extracted from the fitting of experimental data in a). (c, d) the effect of bottle temperature on film growth rate per cycle (c) and refractive index (d) for the films deposited with thermal and plasma-assisted approaches using Er(CpMe)₃ and Er(tmhd)₃ precursors, respectively. RFP = 2000 W (c, d) for Er(tmhd)₃. The different points for the same deposition temperature originate from the different precursor pulse/purge times.</figcaption></figure>

fitting of ellipsometry spectra, typical dispersion law was obtained along with the variation of absorption coefficient (Fig. 2,b). The analysis of spectral dependence of absorption coefficient permits to estimate the bandgap energy (Eg) of Er₂O₃ films. For this purpose, the Tauc relation between the absorption coefficient (α) and the bandgap (Eg) *ahν = A*(hν - Eg)ⁿ, where hν is the photon energy, A is a constant, and m is a constant that depends on the type of transitions involved, i.e. m = 1/2, 3/2, 2 and 3 corresponds to direct allowed, direct forbidden, indirect allowed and indirect forbidden transitions, respectively. It was reported that the energy of direct bandgap may represent the onset of intrinsic absorption of the oxide, while the energy of indirect bandgap may represent the onset of absorption involving some defect states [39]. Thus, the plotting of these dependencies permits to get information on the bandgap value and defects in our materials. The direct energy gap was obtained by plotting (*ahν*)² versus hν (Fig. 5S,a) whereas the indirect energy gap was obtained by plotting (*ahν*)¹/² versus hν (Fig. 5S,b). It turned out that in the first case Eg ~ 5.76 eV and in the latter case -Eg ~ 5.2 eV that is in agreement with the data reported in [10,39–41].

The growth rate per cycle, r<sub>G</sub>, and corresponding n values were analyzed versus different deposition conditions. For all the films, the effects of the substrate temperature, temperature of bottle with Er precursor, and pulse/purge time for Er precursor and oxidant were considered. Along with this, for thermal ALD, the effect of the boost flow of Er precursor was studied, whereas for the plasma-enhanced ALD, the influence of oxygen plasma power (RFP) and duration of its application was investigated. The obtained results are collected in Figs. 1S–4S.

It turned out that for the films grown with plasma-assisted ALD, the start of film growth with the r<sub>G</sub> = 0.5 pm/cycle occurs when T<sub>b</sub> = 155 °C (Fig. 2,c). Further T<sub>b</sub> rise up to 200 °C results in the r<sub>G</sub> rise. Depending on the Er(tmhd)<sub>3</sub> and O<sub>2</sub> pulse/purge durations, the r<sub>G</sub> can reach the values

from 9 to 20 pm/cycle, being higher for shorter Er pulse time (Fig. 1S). The decrease of O₂ pulse time brings the same result due to incomplete oxidation of Er₂ precursor. It is worth to point out that the T<sub>b</sub> = 200 °C value for the Er(tmhd)<sub>3</sub> cylinder exceeds significantly the values reported for similar Er compounds in [12] that could be caused by the difference in the configuration of the cylinders used in [12] and in our ALD unit.

The r<sub>G</sub> value has a non-monotonic behavior versus oxygen plasma power (Fig. 2S). The highest value of about 20 pm/cycle is obtained for RFP =2000 W. This RFP value was kept for further production of the films with plasma-assisted approach.

The refractive index of the films deposited at T<sub>b</sub> = 200 °C is found to be in the range of n =1.75–1.92 showing a trend to decrease with Er pulse time rise (Figs. 2,d, 1S). The obtained values are in good agreement with those reported in the literature for Er₂O₃ [5,40].

For the films grown with thermal ALD approach, no deposition is observed for T<sub>b</sub> < 120 °C. For T<sub>b</sub> = 120 °C, the appearance of the film with r<sub>G</sub> 0.5 pm/cycle is detected (Fig. 2,c). Further T<sub>b</sub> increase up to 170 °C results in the r<sub>G</sub> increasing up to 80–100 pm/cycle whereas at higher T<sub>b</sub>, the r<sub>G</sub> decreases dramatically and for T<sub>b</sub> = 190–200 °C no film is obtained (Figs. 2,c, 3S).

The highest r<sub>G</sub> values of about 80-100 pm/cycle was achieved for the films deposited with T<sub>b</sub> = 170 °C and t<sub>pulse</sub>/t<sub>purge</sub> = (1.6–2.0)s/4s (Figs. 2, c, 3S). The refractive index for these films measured at 1.95 eV is found to be n = 1.8–1.87 being similar to the value observed for films grown with the plasma-assisted approach. It should be noted that further T<sub>b</sub> increase up to 175–180 °C caused the crystallization of Er(CpMe)<sub>3</sub> precursor inside the cylinder and, as a consequence, no film deposition on substrate. The increase of the Er(CpMe)<sub>3</sub> pulse time from 1.6 s up to 2 s results in the slight increase of growth rate, whereas further pulse time
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** (a) Typical ellipsometry spectra for as-deposited Er₂O₃ film (symbols) taken at 65°, 70° and 75° and their fitting (solid curves) using "new amorphous" model. (b) Refractive index and absorption coefficient of as-deposited Er₂O₃ film extracted from the fitting of experimental data in a). (c, d) the effect of bottle temperature on film growth rate per cycle (c) and refractive index (d) for the films deposited with thermal and plasma-assisted approaches using Er(CpMe)₃ and Er(tmhd)₃ precursors, respectively. RFP = 2000 W (c, d) for Er(tmhd)₃. The different points for the same deposition temperature originate from the different precursor pulse/purge times.</figcaption></figure>

fitting of ellipsometry spectra, typical dispersion law was obtained along with the variation of absorption coefficient (Fig. 2,b). The analysis of spectral dependence of absorption coefficient permits to estimate the bandgap energy (E_g) of Er₂O₃ films. For this purpose, the Tauc relation between the absorption coefficient (α) and the bandgap (E_g) *ahν = A *(*hν - E_g*)ⁿ, where hν is the photon energy, A is a constant, and m is a constant that depends on the type of transitions involved, i.e. m = 1/2, 3/2, 2 and 3 corresponds to direct allowed, direct forbidden, indirect allowed and indirect forbidden transitions, respectively. It was reported that the energy of direct bandgap may represent the onset of intrinsic absorption of the oxide, while the energy of indirect bandgap may represent the onset of absorption involving some defect states [39]. Thus, the plotting of these dependencies permits to get information on the bandgap value and defects in our materials. The direct energy gap was obtained by plotting (*ahν*)² versus hν (Fig. 5S,a) whereas the indirect energy gap was obtained by plotting (*ahν*)¹/² versus hν (Fig. 5S,b). It turned out that in the first case E_g ~ 5.76 eV and in the latter case -E_g ~ 5.2 eV that is in agreement with the data reported in [10,39–41].

The growth rate per cycle, r<sub>G</sub>, and corresponding n values were analyzed versus different deposition conditions. For all the films, the effects of the substrate temperature, temperature of bottle with Er precursor, and pulse/purge time for Er precursor and oxidant were considered. Along with this, for thermal ALD, the effect of the boost flow of Er precursor was studied, whereas for the plasma-enhanced ALD, the influence of oxygen plasma power (RFP) and duration of its application was investigated. The obtained results are collected in Figs. 1S–4S.

It turned out that for the films grown with plasma-assisted ALD, the start of film growth with the r<sub>G</sub> = 0.5 pm/cycle occurs when T<sub>b</sub> = 155 °C (Fig. 2,c). Further T<sub>b</sub> rise up to 200 °C results in the r<sub>G</sub> rise. Depending on the Er(tmhd)<sub>3</sub> and O<sub>2</sub> pulse/purge durations, the r<sub>G</sub> can reach the values

from 9 to 20 pm/cycle, being higher for shorter Er pulse time (Fig. 1S). The decrease of O₂ pulse time brings the same result due to incomplete oxidation of Er₂ precursor. It is worth to point out that the T<sub>b</sub> = 200 °C value for the Er(tmhd)<sub>3</sub> cylinder exceeds significantly the values reported for similar Er compounds in [12] that could be caused by the difference in the configuration of the cylinders used in [12] and in our ALD unit.

The r<sub>G</sub> value has a non-monotonic behavior versus oxygen plasma power (Fig. 2S). The highest value of about 20 pm/cycle is obtained for RFP =2000 W. This RFP value was kept for further production of the films with plasma-assisted approach.

The refractive index of the films deposited at T<sub>b</sub> = 200 °C is found to be in the range of n =1.75–1.92 showing a trend to decrease with Er pulse time rise (Figs. 2,d, 1S). The obtained values are in good agreement with those reported in the literature for Er₂O₃ [5,40].

For the films grown with thermal ALD approach, no deposition is observed for T<sub>b</sub> < 120 °C. For T<sub>b</sub> = 120 °C, the appearance of the film with r<sub>G</sub> 0.5 pm/cycle is detected (Fig. 2,c). Further T<sub>b</sub> increase up to 170 °C results in the r<sub>G</sub> increasing up to 80–100 pm/cycle whereas at higher T<sub>b</sub>, the r<sub>G</sub> decreases dramatically and for T<sub>b</sub> = 190–200 °C no film is obtained (Figs. 2,c, 3S).

The highest r<sub>G</sub> values of about 80-100 pm/cycle was achieved for the films deposited with T<sub>b</sub> = 170 °C and t<sub>pulse</sub>/t<sub>purge</sub> = (1.6–2.0)s/4s (Figs. 2, c, 3S). The refractive index for these films measured at 1.95 eV is found to be n = 1.8–1.87 being similar to the value observed for films grown with the plasma-assisted approach. It should be noted that further T<sub>b</sub> increase up to 175–180 °C caused the crystallization of Er(CpMe)<sub>3</sub> precursor inside the cylinder and, as a consequence, no film deposition on substrate. The increase of the Er(CpMe)<sub>3</sub> pulse time from 1.6 s up to 2 s results in the slight increase of growth rate, whereas further pulse time
<figure><img src="image_4.png" /><figcaption>Fig. 3. (a) Evolution of the refractive index (at 1.95 eV) and (b) thickness of the films grown with the thermal approach (square symbols) and the plasma-assisted approach (circle symbols) with annealing temperature $T_A$. The symbols marked by "AD" and dashed lines at $T = 300$ °C correspond to the thickness and refractive index of as-deposited films that have been grown at 300 °C.</figcaption></figure>

rise up to 4 s did not change it (Fig. 3S)

Analysis of the effect of deposition conditions reveals a low refractive index ($n = 1.65$) for the films grown from $Er(CpMe)_3$ with $T_b = 120$ °C and from $Er(tmhd)_3$ with $T_b = 155$ °C. The lowering of the $n$ value can be caused by the presence of a phase with lower refractive index in the film due to incomplete Er precursor oxidation and/or presence of byproducts, as well as the formation of $SiO_x$ or $ErSiO_x$ layer on the substrate surface instead of $Er_2O_3$ film in contrast to the results reported in [12] where stoichiometric $Er_2O_3$ films were obtained.

The effect of substrate temperature on the growth rate and optical properties of the films was also studied. The optimal deposition temperature is determined as $T_d=300$°C which allowed achieving high growth rate for both approaches. Taking into account the results described above, the effect of annealing treatment on the properties of the films grown with thermal ALD approach (with $T_d = 300$ °C, $T_b = 170$ °C, $Er(CpMe)_3/H_2O$ pulse/purge time as 1.6s/2.0s/0.1s/2.0s) and by plasma-assisted deposition (with $T_d = 300$ °C, $T_b = 200$ °C, $Er(tmhd)_3/O_2$ plasma pulse/purge time as 2.0s/4.0s/26.0s/4.0s) will be discussed in details.

## 3.2. Effect of annealing treatment on film thickness and refractive index

Fig. 3 represents the effect of annealing temperature on the refractive index at 1.95 eV (Fig. 3a) and thickness (Fig. 3b) of the films produced with the two different approaches. It is seen that, the refractive index of as-deposited films grown with the plasma-assisted approach is about $n = 1.860$ (at 300 °C). As evidenced, the annealing treatment does not modify greatly the values of the refractive index as well as the thickness of the film. By contrast, annealing at $T_A = 500$ °C results in the slight increase of the $n$ up to 1.896 that remains almost constant with $T_A$ rise up to 800 °C. For $T_A = 900$ °C, the refractive index reaches $n = 1.920$, while further $T_A$ rise up to 1000 °C causes some decrease of $n$ down to 1.90. At the same time, the thickness of the film decreases slightly when $T_A = 500$ °C and does not change up to $T_A = 900$ °C. Such a behavior can be ascribed to the densification of the film and its crystallization. For $T_A = 1000$ °C, the film thickness becomes larger, but does not exceed the thickness of the as-deposited film. Taking into account some decrease of the refractive index along with the film thickening for $T_A = 1000$ °C, one can suppose that the formation of a phase with lower refractive index (for instance, $SiO_x$ or $ErSiO_x$ as well as cavities) can occur upon

annealing.

The films, prepared with the thermal approach, demonstrate pronounced variation of the refractive index and thickness with annealing treatment. With $T_A$ rise up to 800 °C, the $n$ value evolves from 1.84 (as-deposited at 300°C) up to 1.95 ($T_A = 800$ °C). Simultaneously, the thickness of the film increases from 108 up to 115 nm ($T_A = 500$ °C) and does not change up to $T_A = 800-900$ °C. The increase of the refractive index with a concomitant unchangeable thickness of the film can be ascribed to the film crystallization. However, further $T_A$ increase up to 1000 °C results in the decrease of refractive index down to $n = 1.83$ accompanied by the decrease of film thickness. This behavior can be caused by the formation of cavities along with an increasing $SiO_2$ interfacial layer thickness as well as a possible appearance of Er silicate phases.

To obtain more information about structural variation and possible formation of pores and/or Si-rich phases, the samples were studied with FTIR and XRD methods as well as different electron microscopy techniques.

## 3.3. Effect of annealing on structural properties of the films

### 3.3.1. AFM data

All the films described above were observed using optical and atomic force microscopy methods. Fig. 4,a-c show the surface morphology of $Er_2O_3$ films grown with the thermal approach and annealed at different temperatures. It is seen that as-deposited film has a smooth surface that is free of any structural elements (Fig. 4,a).

Annealing results in the appearance of some contrasts on the film surface as well as formation of holes and/or cracks (Fig. 4,b and c). These latter are observed in small quantities for the films heated at $T_A = 600$ °C (Fig. 4,b), while annealing at $T_A = 900$ °C results in the formation of a crack network (Fig. 4,c). Further $T_A$ increase (up to 1000 °C) causes the development of the holes and cracks as well as the increase of their depth. This structure evolution was found to be similar to the data reported in Ref. [41].

The surface of the samples prepared with plasma assisted approach is found to be smooth both for as-deposited samples and those annealed at $T_A \le 800$ °C (Fig. 4,d and f). No cracks were found in such films whatever the $T_A$ value. For $T_A = 900-1000$ °C, a modification of the surface morphology occurs (Fig. 4, f). However, these transformations are less significant than for the films grown with the thermal approach. These findings are confirmed by observations of the same samples with AFM method and presented hereafter.

Fig. 5 represents AFM images of the surface of the films annealed at high temperatures ($T_A = 900$ and $1000$ °C). It is seen that the surface of the films, grown with the plasma-assisted approach and annealed at $900$ °C, presents structural elements (or grains).

The analysis of AFM image over $2 \times 2 \mu m^2$ surface (Fig. 5,a) shows that these grains have lateral and vertical sizes in the range 20–140 nm and 3–10 nm, respectively. The root-mean square roughness (RMS) is about 1.03 nm. Annealing at $T_A = 1000$ °C results in the RMS increase up to 1.39 nm. The lateral sizes of the elements (grains) rise up to 50–170 nm, whereas their height is still in the 5–10 nm range (Fig. 5,b).

AFM study of the films grown with thermal ALD approach revealed also that heat treatment caused not only the appearance of the grains on film surface, but also the development of a crack network (Fig. 5,c). Indeed, for the films annealed at $T_A = 900$ °C, for the observed regions, grains with lateral sizes of 20–200 nm and height of 3–8 nm are detected, whereas the width of the cracks is 25–100 nm. As for the crack depth, its value is found to be between 10 and 25 nm. However, one can note that the cracks could be deeper, but AFM tip cannot penetrate in the crack depth due to its physical size. The RMS value was 2.61 nm (over $2 \times 2 \mu m^2$ surface), whereas the RMS value of the plateau between the cracks was about 1.25 nm (Fig. 5,c).

The heating of the samples at $T_A = 1000$ °C induced also the development of the cracks (Fig. 5,d). Along with this, the number of
<figure><img src="image_5.png" /><figcaption>Fig. 3. (a) Evolution of the refractive index (at 1.95 eV) and (b) thickness of the films grown with the thermal approach (square symbols) and the plasma-assisted approach (circle symbols) with annealing temperature $T_A$. The symbols marked by "AD" and dashed lines at $T = 300$ °C correspond to the thickness and refractive index of as-deposited films that have been grown at 300 °C.</figcaption></figure>

rise up to 4 s did not change it (Fig. 3S)

Analysis of the effect of deposition conditions reveals a low refractive index ($n = 1.65$) for the films grown from $Er(CpMe)_3$ with $T_b = 120$ °C and from $Er(tmhd)_3$ with $T_b = 155$ °C. The lowering of the $n$ value can be caused by the presence of a phase with lower refractive index in the film due to incomplete Er precursor oxidation and/or presence of byproducts, as well as the formation of $SiO_x$ or $ErSiO_x$ layer on the substrate surface instead of $Er_2O_3$ film in contrast to the results reported in [12] where stoichiometric $Er_2O_3$ films were obtained.

The effect of substrate temperature on the growth rate and optical properties of the films was also studied. The optimal deposition temperature is determined as $T_d=300$°C which allowed achieving high growth rate for both approaches. Taking into account the results described above, the effect of annealing treatment on the properties of the films grown with thermal ALD approach (with $T_d = 300$ °C, $T_b = 170$ °C, $Er(CpMe)_3/H_2O$ pulse/purge time as 1.6s/2.0s/0.1s/2.0s) and by plasma-assisted deposition (with $T_d = 300$ °C, $T_b = 200$ °C, $Er(tmhd)_3/O_2$ plasma pulse/purge time as 2.0s/4.0s/26.0s/4.0s) will be discussed in details.

## 3.2. Effect of annealing treatment on film thickness and refractive index

Fig. 3 represents the effect of annealing temperature on the refractive index at 1.95 eV (Fig. 3a) and thickness (Fig. 3b) of the films produced with the two different approaches. It is seen that, the refractive index of as-deposited films grown with the plasma-assisted approach is about $n = 1.860$ (at 300 °C). As evidenced, the annealing treatment does not modify greatly the values of the refractive index as well as the thickness of the film. By contrast, annealing at $T_A = 500$ °C results in the slight increase of the $n$ up to 1.896 that remains almost constant with $T_A$ rise up to 800 °C. For $T_A = 900$ °C, the refractive index reaches $n = 1.920$, while further $T_A$ rise up to 1000 °C causes some decrease of $n$ down to 1.90. At the same time, the thickness of the film decreases slightly when $T_A = 500$ °C and does not change up to $T_A = 900$ °C. Such a behavior can be ascribed to the densification of the film and its crystallization. For $T_A = 1000$ °C, the film thickness becomes larger, but does not exceed the thickness of the as-deposited film. Taking into account some decrease of the refractive index along with the film thickening for $T_A = 1000$ °C, one can suppose that the formation of a phase with lower refractive index (for instance, $SiO_x$ or $ErSiO_x$ as well as cavities) can occur upon

annealing.

The films, prepared with the thermal approach, demonstrate pronounced variation of the refractive index and thickness with annealing treatment. With $T_A$ rise up to 800 °C, the $n$ value evolves from 1.84 (as-deposited at 300°C) up to 1.95 ($T_A = 800$ °C). Simultaneously, the thickness of the film increases from 108 up to 115 nm ($T_A = 500$ °C) and does not change up to $T_A = 800–900$ °C. The increase of the refractive index with a concomitant unchangeable thickness of the film can be ascribed to the film crystallization. However, further $T_A$ increase up to 1000 °C results in the decrease of refractive index down to $n = 1.83$ accompanied by the decrease of film thickness. This behavior can be caused by the formation of cavities along with an increasing $SiO_2$ interfacial layer thickness as well as a possible appearance of Er silicate phases.

To obtain more information about structural variation and possible formation of pores and/or Si-rich phases, the samples were studied with FTIR and XRD methods as well as different electron microscopy techniques.

## 3.3. Effect of annealing on structural properties of the films

### 3.3.1. AFM data

All the films described above were observed using optical and atomic force microscopy methods. Fig. 4,a-c show the surface morphology of $Er_2O_3$ films grown with the thermal approach and annealed at different temperatures. It is seen that as-deposited film has a smooth surface that is free of any structural elements (Fig. 4,a).

Annealing results in the appearance of some contrasts on the film surface as well as formation of holes and/or cracks (Fig. 4,b and c). These latter are observed in small quantities for the films heated at $T_A = 600$ °C (Fig. 4,b), while annealing at $T_A = 900$ °C results in the formation of a crack network (Fig. 4,c). Further $T_A$ increase (up to 1000 °C) causes the development of the holes and cracks as well as the increase of their depth. This structure evolution was found to be similar to the data reported in Ref. [41].

The surface of the samples prepared with plasma assisted approach is found to be smooth both for as-deposited samples and those annealed at $T_A \le 800$ °C (Fig. 4,d and f). No cracks were found in such films whatever the $T_A$ value. For $T_A = 900–1000$ °C, a modification of the surface morphology occurs (Fig. 4, f). However, these transformations are less significant than for the films grown with the thermal approach. These findings are confirmed by observations of the same samples with AFM method and presented hereafter.

Fig. 5 represents AFM images of the surface of the films annealed at high temperatures ($T_A = 900$ and $1000$ °C). It is seen that the surface of the films, grown with the plasma-assisted approach and annealed at $900$ °C, presents structural elements (or grains).

The analysis of AFM image over $2 \times 2 \mu m^2$ surface (Fig. 5,a) shows that these grains have lateral and vertical sizes in the range 20–140 nm and 3–10 nm, respectively. The root-mean square roughness (RMS) is about 1.03 nm. Annealing at $T_A = 1000$ °C results in the RMS increase up to 1.39 nm. The lateral sizes of the elements (grains) rise up to 50–170 nm, whereas their height is still in the 5–10 nm range (Fig. 5,b).

AFM study of the films grown with thermal ALD approach revealed also that heat treatment caused not only the appearance of the grains on film surface, but also the development of a crack network (Fig. 5,c). Indeed, for the films annealed at $T_A = 900$ °C, for the observed regions, grains with lateral sizes of 20–200 nm and height of 3–8 nm are detected, whereas the width of the cracks is 25–100 nm. As for the crack depth, its value is found to be between 10 and 25 nm. However, one can note that the cracks could be deeper, but AFM tip cannot penetrate in the crack depth due to its physical size. The RMS value was 2.61 nm (over $2 \times 2 \mu m^2$ surface), whereas the RMS value of the plateau between the cracks was about 1.25 nm (Fig. 5,c).

The heating of the samples at $T_A = 1000$ °C induced also the development of the cracks (Fig. 5,d). Along with this, the number of
<figure><img src="image_5.png" /><figcaption>**Fig. 4.** Surface images obtained with an optical microscope for as-deposited and annealed Er₂O₃ films grown with thermal and plasma-assisted approaches using Er (CpMe)₃ (a–c) and Er(tmhd)₃ (d–f) precursors, respectively. Annealing temperatures are mentioned in the figures.</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>**Fig. 5.** AFM images of the surface of Er₂O₃ films grown with the plasma assisted approach (a, b) and the thermal approach (c, d) and annealed at 900 °C (a, c) and 1000 °C (b, d).</figcaption></figure>

grains on the plateau between the cracks increases. Two types of the grains can be seen on the top surface of the plateau, i.e. small gains with average size of about 100–140 nm and large grains with the lateral sizes

of 200–250 nm. It would be noted that the measured depth of the cracks' could be underestimated because of the curvature radius of the AFM tip. TEM results presented in the following sections, support this statement
<figure><img src="image_6.png" /><figcaption>**Fig. 4.** Surface images obtained with an optical microscope for as-deposited and annealed Er₂O₃ films grown with thermal and plasma-assisted approaches using Er (CpMe)₃ (a–c) and Er(tmhd)₃ (d–f) precursors, respectively. Annealing temperatures are mentioned in the figures.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Fig. 5.** AFM images of the surface of Er₂O₃ films grown with the plasma assisted approach (a, b) and the thermal approach (c, d) and annealed at 900 °C (a, c) and 1000 °C (b, d).</figcaption></figure>

grains on the plateau between the cracks increases. Two types of the grains can be seen on the top surface of the plateau, i.e. small gains with average size of about 100–140 nm and large grains with the lateral sizes

of 200–250 nm. It would be noted that the measured depth of the cracks' could be underestimated because of the curvature radius of the AFM tip. TEM results presented in the following sections, support this statement
and show that in some regions the cracks penetrate down to the sub-
strate surface (Fig. 6S, see supplement materials).

The results presented above show that the films grown with the plasma-assisted approach are structurally more stable upon anneal treatment. However, in both cases, the formation of grains on the film surface can be the result of its crystallization. Due to mismatching in lattice parameters between Er₂O₃ film (for cubic Er₂O₃, a =1.054 nm) and underlying Si substrate (for cubic Si, a =0.543 nm), one can expect that films are compressively stressed. Thus, an annealing can result in their relaxation via the formation some cracks. However, in the films grown with the plasma assisted approach, such stresses can be smaller than in their counterparts prepared with the thermal approach. One of the reasons could be the higher growth rate per cycle in the thermal approach (about 80–100 pm/cycle against 15–20 pm/cycle for plasma assisted approach). The presence of OH- or CH-groups as well as other volatile complexes (that were not completely removed during the deposition) can be an additional reason of this phenomenon. Thermal treatment at high T<sub>A</sub> values can also provoke the diffusion and desorp- tion of such complexes and, as a consequence, the formation of cracks or even film delamination from the substrate. Besides, the formation of silicates allowing a decrease in the mismatching between the film and underlying substrate is also possible. One can assume that the mecha- nism of the film relaxation could be different for the films grown with different approaches.

3.3.2. XRD data

In order to see the effect of annealing on the film crystallization, XRD
study was performed for the samples mentioned above and the corre-
sponding patterns are plotted in Fig. 6. The comparison of XRD data with
the standard cubic Er₂O₃ patterns [29,30] shows that the films are
crystallized in cubic phase, even after deposition. Both grazing and
Bragg-Brentano geometries support this statement (Fig. 6,a and b,
respectively). As-deposited films grown at 300 °C show the most intense

peaks in the 2θ = 20-60° range at 2θ = 23.48° (2 1 1), 33.73° (2 2 2), 39.10° (4 0 0) and 56.82° (4 4 0) (Fig. 6), that are shifted towards lower angles in comparison with standard values reported for cubic Er₂O₃. This shift is most probably due to the mismatch between the lattice parameters of the film and the underlying substrate. Indeed, for standard Er₂O₃, the a<sub>Er2O3</sub> = 1.054 nm value is about twice larger than the unit cell parameter of Si substrate (2a<sub>Si</sub> = 1.086 nm). However, a small negative mismatch for the Er₂O₃ film can generate stress. Thus, for as-deposited films, it is found that a<sub>Er2O3</sub> = 1.0617 nm and the micro-strains are about ~0.74%.

Thermal treatment of the films at high temperatures causes further
their crystallization as well as strain relaxation. Indeed, the T<sub>A</sub> increase
up to 1000°C stimulates some narrowing of XRD peaks and their gradual
shift towards higher angles, i.e. 2θ = 24.06° (2 1 1), 34.12° (2 2 2),
39.64° (4 0 0) and 57.36° (4 4 0) (Fig. 6,b), approaching the values of
standard cubic Er<sub>2</sub>O<sub>3</sub>. The highest intensity of the peak at 2θ = 33.73° (2
2 2) is the evidence of the preferred orientation of the film following the
[111] direction. The analysis of XRD data revealed the gradual decrease
of the lattice parameters down to 1.057 nm and the increase of coherent
domain size from ~9 nm up to ~14 nm as well as stress relaxation down
to 0.17%. No any other Er<sub>2</sub>O<sub>3</sub> phases are identified in the films annealed
at T<sub>A</sub> ≤ 1000 °C. For T<sub>A</sub> = 1100 °C, additional XRD peaks appear at 2θ =
35.36° and 38.15° that can be assigned to Er<sub>2</sub>SiO<sub>5</sub> phase (PDF card No.
52-1809). An appearance of this phase is caused by the diffusion of
silicon atoms from underling substrate in the film volume due to the
high temperature applied. The appearance of silicate phase was also
reported for Er<sub>2</sub>O<sub>3</sub> thin films by other groups [19,20],.

Thus, XRD data confirm the conclusions drawn from ellipsometry
data and AFM observations. However, to obtain more information about
the factors that can result in the evolution of the refractive index and of
the film thicknesses, FTIR experiments were carried out. Besides, this
method permits to detect the OH-groups (if any) and the formation of
SiO<sub>2</sub> or Er-silicate phases in the films described above.

<figure><img src="image_7.png" /><figcaption>**Fig. 6.** Grazing-incidence XRD patterns of Er₂O₃ films grown with plasma-assisted (a) and thermal (b) approaches. The conditions of XRD experiment: (a) ω = 0.25, 0.5 and 1.05°; excitation source – Cu Kα = 1.5406 Å, the patterns are recorded for films annealed at 500 °C; (b) ω = 0.5°; excitation source – Co Kα = 1.78897 Å, XRD patterns are recorded for as-deposited and annealed at T<sub>A</sub> = 800–1000 °C films. In both figures, the curves are shifted vertically for clarity.</figcaption></figure>
and show that in some regions the cracks penetrate down to the sub-
strate surface (Fig. 6S, see supplement materials).

The results presented above show that the films grown with the
plasma-assisted approach are structurally more stable upon anneal
treatment. However, in both cases, the formation of grains on the film
surface can be the result of its crystallization. Due to mismatching in
lattice parameters between Er₂O₃ film (for cubic Er₂O₃, a =1.054 nm)
and underlying Si substrate (for cubic Si, a =0.543 nm), one can expect
that films are compressively stressed. Thus, an annealing can result in
their relaxation via the formation some cracks. However, in the films
grown with the plasma assisted approach, such stresses can be smaller
than in their counterparts prepared with the thermal approach. One of
the reasons could be the higher growth rate per cycle in the thermal
approach (about 80–100 pm/cycle against 15–20 pm/cycle for plasma
assisted approach). The presence of OH- or CH-groups as well as other
volatile complexes (that were not completely removed during the
deposition) can be an additional reason of this phenomenon. Thermal
treatment at high T<sub>A</sub> values can also provoke the diffusion and desorp-
tion of such complexes and, as a consequence, the formation of cracks or
even film delamination from the substrate. Besides, the formation of
silicates allowing a decrease in the mismatching between the film and
underlying substrate is also possible. One can assume that the mecha-
nism of the film relaxation could be different for the films grown with
different approaches.

3.3.2. XRD data

In order to see the effect of annealing on the film crystallization, XRD
study was performed for the samples mentioned above and the corre-
sponding patterns are plotted in Fig. 6. The comparison of XRD data with
the standard cubic Er₂O₃ patterns [29,30] shows that the films are
crystallized in cubic phase, even after deposition. Both grazing and
Bragg-Brentano geometries support this statement (Fig. 6,a and b,
respectively). As-deposited films grown at 300 °C show the most intense

peaks in the 2θ = 20-60° range at 2θ = 23.48° (2 1 1), 33.73° (2 2 2), 39.10° (4 0 0) and 56.82° (4 4 0) (Fig. 6), that are shifted towards lower angles in comparison with standard values reported for cubic Er₂O₃. This shift is most probably due to the mismatch between the lattice parameters of the film and the underlying substrate. Indeed, for standard Er₂O₃, the a<sub>Er2O3</sub> = 1.054 nm value is about twice larger than the unit cell parameter of Si substrate (2a<sub>Si</sub> = 1.086 nm). However, a small negative mismatch for the Er₂O₃ film can generate stress. Thus, for as-deposited films, it is found that a<sub>Er2O3</sub> = 1.0617 nm and the micro-strains are about ~0.74%.

Thermal treatment of the films at high temperatures causes further
their crystallization as well as strain relaxation. Indeed, the T<sub>A</sub> increase
up to 1000°C stimulates some narrowing of XRD peaks and their gradual
shift towards higher angles, i.e. 2θ = 24.06° (2 1 1), 34.12° (2 2 2),
39.64° (4 0 0) and 57.36° (4 4 0) (Fig. 6,b), approaching the values of
standard cubic Er<sub>2</sub>O<sub>3</sub>. The highest intensity of the peak at 2θ = 33.73° (2
2 2) is the evidence of the preferred orientation of the film following the
[111] direction. The analysis of XRD data revealed the gradual decrease
of the lattice parameters down to 1.057 nm and the increase of coherent
domain size from ~9 nm up to ~14 nm as well as stress relaxation down
to 0.17%. No any other Er<sub>2</sub>O<sub>3</sub> phases are identified in the films annealed
at T<sub>A</sub> ≤ 1000 °C. For T<sub>A</sub> = 1100 °C, additional XRD peaks appear at 2θ =
35.36° and 38.15° that can be assigned to Er<sub>2</sub>SiO<sub>5</sub> phase (PDF card No.
52-1809). An appearance of this phase is caused by the diffusion of
silicon atoms from underling substrate in the film volume due to the
high temperature applied. The appearance of silicate phase was also
reported for Er<sub>2</sub>O<sub>3</sub> thin films by other groups [19,20],.

Thus, XRD data confirm the conclusions drawn from ellipsometry
data and AFM observations. However, to obtain more information about
the factors that can result in the evolution of the refractive index and of
the film thicknesses, FTIR experiments were carried out. Besides, this
method permits to detect the OH-groups (if any) and the formation of
SiO<sub>2</sub> or Er-silicate phases in the films described above.

<figure><img src="image_8.png" /><figcaption>Fig. 6. Grazing-incidence XRD patterns of Er₂O₃ films grown with plasma-assisted (a) and thermal (b) approaches. The conditions of XRD experiment: (a) ω = 0.25, 0.5 and 1.05°; excitation source – Cu Kα = 1.5406 Å, the patterns are recorded for films annealed at 500 °C; (b) ω = 0.5°; excitation source – Co Kα = 1.78897 Å, XRD patterns are recorded for as-deposited and annealed at T<sub>A</sub> = 800–1000 °C films. In both figures, the curves are shifted vertically for clarity.</figcaption></figure>
### 3.3.3. Evolution of FTIR spectra

The FTIR spectra of the as-deposited Er₂O₃ film grown different ALD approaches were found to be similar that is an evidence of their microstructural similarity. The detailed analysis of FTIR spectra recorded for mid-IR spectral range for the samples grown with Er(CpMe)₃ precursor was performed by us earlier in Ref. [42]. Hereafter, the spectra of the films grown with plasma-assisted ALD approach are presented.

Fig. 7 shows FTIR spectra of Er₂O₃ films grown with Er(tmhd)₃ precursor detected in mid-IR spectral range under normal and 60° incidences of excited light. The signals of Er-O vibrations in the 400–650 cm⁻¹ range are seen (Fig. 7,a). Besides, the modes caused by C-H vibrations as well as H₂O and OH-related ones are present (Fig. 7,b). These vibration bands could originate from the byproducts remained in the films after precursor decomposition and oxidation. An annealing resulted in the decrease of these signals up to their vanishing for Tₐ > 900 °C.

Apart from this, in the films annealed at T<sub>A</sub> ≥700°C, Er-O related peaks become to be well-defined, being sharper for higher T<sub>A</sub> values (Fig. 7). Furthermore, additional Si-O-Si bands peaked at about 1250, 1080 and 820 cm<sup>-1</sup> appear. Since the films were grown on Si substrate covered with thermal 5-nm SiO<sub>2</sub> layer, the increase of the magnitude of the Si-O-Si band peaked at 1250 cm<sup>-1</sup> is the evidence of the increase of SiO<sub>2</sub> layer thickness that can be caused by the oxygen diffusion from the film towards film/Si substrate interface. Similar phenomenon was reported in [[36,43]. It was also shown that oxygen diffusion is 5–6 orders of magnitude faster along the grain boundaries than in the bulk [44] that can explain the increase of SiO<sub>x</sub> layer thickness. Further T<sub>A</sub> increase results in the decrease of the contribution of Si-O-Si related bands as well as in the appearance and enhancement of Si-O-Er bands in the 800–1100 cm<sup>-1</sup> spectral range. The presence of the Si-O-Er bands could originate from the formation Er silicate phase due to interaction of the film with underlying substrate via Er diffusion towards film/substrate interface and/or Si diffusion in film volume.

Similar morphology evolution for the films grown with two approaches is demonstrated by Fig. 8 where FTIR spectra are compared for

mid-IR and far-IR spectral ranges. It is seen that as-deposited films of both types show Er-O related vibration bands in the 100–700 cm⁻¹ range. Their narrowing occurred upon film annealing at T<sub>A</sub> ≥ 900 °C indicating crystallization of the films which is in agreement with the XRD results. At the same time, the comparison of the shape of Si-O-Er related peaks revealed that these peaks are sharper for the films grown with plasma-assisted ALD. Thus one can assume that more pronounced crystallization occurs in the films grown with Er(tmhd)₃ precursor.

The formation of Er silicate phase was reported in Refs. [39,43,45]. It was shown that this process can require a short annealing time, for instance, it can occur upon rapid thermal annealing (at 600–900 °C for 30 s). As for the Er₂O₃ films studied in this work, they are grown as crystalline. The presence of grain boundaries can be a reason not only for fast oxygen diffusion from the film towards film/substrate interface upon annealing, but also for the diffusion of other species, for instance, Si atoms from the substrate (or interface layer) in the film volume.

Since in our case annealing time was longer (30 min) the diffusion process can be very significant. Besides, high-temperature annealing stimulates further film crystallization that can promote such diffusion via the appearance of additional grain boundaries. Meanwhile, the broad featureless Si-O-Er vibration band as well as nearly unchangeable intensity of Si-O-Si vibration band peaked at 1250 cm⁻¹ detected in the films grown with Er(CpMe)₃ and annealed at T<sub>A</sub> ≤ 1000 °C could be caused by the difference in the diffusion processes and film crystallization in comparison with similar processes occurring in the films grown with Er(tmhd)₃ precursor and submitted to the same thermal treatment. The higher ratio of the magnitudes of Er-O bands to that of Si-O-Si vibrations observed for the films grown with Er(CpMe)₃ demonstrates the better crystallization of Er₂O₃ original layer and lower contribution of Er silicate crystallization that can be caused by the competition of the oxygen diffusion towards film/substrate interface and out-diffusion of byproducts (for instance, OH-groups) whose presence was detected in as-grown films [42]. The crystallization of Er silicate phase was revealed

<figure><img src="image_8.png" /><figcaption>Fig. 7. FTIR spectra of Er₂O₃ films grown with plasma assisted ALD versus Tₐ recorded with normal and 60° incidence of excitation light and presented for different spectral ranges.</figcaption></figure>
### 3.3.3. Evolution of FTIR spectra

The FTIR spectra of the as-deposited Er₂O₃ film grown different ALD approaches were found to be similar that is an evidence of their microstructural similarity. The detailed analysis of FTIR spectra recorded for mid-IR spectral range for the samples grown with Er(CpMe)₃ precursor was performed by us earlier in Ref. [42]. Hereafter, the spectra of the films grown with plasma-assisted ALD approach are presented.

Fig. 7 shows FTIR spectra of Er₂O₃ films grown with Er(tmhd)₃ precursor detected in mid-IR spectral range under normal and 60° incidences of excited light. The signals of Er-O vibrations in the 400–650 cm⁻¹ range are seen (Fig. 7,a). Besides, the modes caused by C-H vibrations as well as H₂O and OH-related ones are present (Fig. 7,b). These vibration bands could originate from the byproducts remained in the films after precursor decomposition and oxidation. An annealing resulted in the decrease of these signals up to their vanishing for Tₐ > 900 °C.

Apart from this, in the films annealed at T<sub>A</sub> ≥700°C, Er-O related peaks become to be well-defined, being sharper for higher T<sub>A</sub> values (Fig. 7). Furthermore, additional Si-O-Si bands peaked at about 1250, 1080 and 820 cm<sup>-1</sup> appear. Since the films were grown on Si substrate covered with thermal 5-nm SiO<sub>2</sub> layer, the increase of the magnitude of the Si-O-Si band peaked at 1250 cm<sup>-1</sup> is the evidence of the increase of SiO<sub>2</sub> layer thickness that can be caused by the oxygen diffusion from the film towards film/Si substrate interface. Similar phenomenon was reported in [[36,43]. It was also shown that oxygen diffusion is 5–6 orders of magnitude faster along the grain boundaries than in the bulk [44] that can explain the increase of SiO<sub>x</sub> layer thickness. Further T<sub>A</sub> increase results in the decrease of the contribution of Si-O-Si related bands as well as in the appearance and enhancement of Si-O-Er bands in the 800–1100 cm<sup>-1</sup> spectral range. The presence of the Si-O-Er bands could originate from the formation Er silicate phase due to interaction of the film with underlying substrate via Er diffusion towards film/substrate interface and/or Si diffusion in film volume.

Similar morphology evolution for the films grown with two approaches is demonstrated by Fig. 8 where FTIR spectra are compared for

mid-IR and far-IR spectral ranges. It is seen that as-deposited films of both types show Er-O related vibration bands in the 100–700 cm⁻¹ range. Their narrowing occurred upon film annealing at T<sub>A</sub> ≥ 900 °C indicating crystallization of the films which is in agreement with the XRD results. At the same time, the comparison of the shape of Si-O-Er related peaks revealed that these peaks are sharper for the films grown with plasma-assisted ALD. Thus one can assume that more pronounced crystallization occurs in the films grown with Er(tmhd)₃ precursor.

The formation of Er silicate phase was reported in Refs. [39,43,45]. It was shown that this process can require a short annealing time, for instance, it can occur upon rapid thermal annealing (at 600–900 °C for 30 s). As for the Er₂O₃ films studied in this work, they are grown as crystalline. The presence of grain boundaries can be a reason not only for fast oxygen diffusion from the film towards film/substrate interface upon annealing, but also for the diffusion of other species, for instance, Si atoms from the substrate (or interface layer) in the film volume.

Since in our case annealing time was longer (30 min) the diffusion process can be very significant. Besides, high-temperature annealing stimulates further film crystallization that can promote such diffusion via the appearance of additional grain boundaries. Meanwhile, the broad featureless Si-O-Er vibration band as well as nearly unchangeable intensity of Si-O-Si vibration band peaked at 1250 cm⁻¹ detected in the films grown with Er(CpMe)₃ and annealed at T<sub>A</sub> ≤ 1000 °C could be caused by the difference in the diffusion processes and film crystallization in comparison with similar processes occurring in the films grown with Er(tmhd)₃ precursor and submitted to the same thermal treatment. The higher ratio of the magnitudes of Er-O bands to that of Si-O-Si vibrations observed for the films grown with Er(CpMe)₃ demonstrates the better crystallization of Er₂O₃ original layer and lower contribution of Er silicate crystallization that can be caused by the competition of the oxygen diffusion towards film/substrate interface and out-diffusion of byproducts (for instance, OH-groups) whose presence was detected in as-grown films [42]. The crystallization of Er silicate phase was revealed

<figure><img src="image_9.png" /><figcaption>Fig. 7. FTIR spectra of Er₂O₃ films grown with plasma assisted ALD versus Tₐ recorded with normal and 60° incidence of excitation light and presented for different spectral ranges.</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Fig. 8. Comparison of FTIR spectra of as-deposited (AD) and annealed at different temperatures Er₂O₃ films grown with Er(CpMe)₃ by the thermal ALD and with Er(tmhd)₃ by the plasma assisted ALD. The spectra were recorded with normal and 60° incidence of excitation light and were shifted vertically in the figures for the clarity. Annealing temperatures are mentioned in the figures.</figcaption></figure>

for the films annealed at $T_A$ = 1100 °C only whereas such process became to be significant already for the films grown with Er(tmhd)₃ precursor at $T_A$ = 900 °C (Figs. 7 and 8).

An analysis of these samples using the TEM technique is necessary to get insight on the spatial localization of SiOₓ and/or Er silicate phases in the samples annealed at high $T_A$. Earlier we have observed cross-sections of the films grown with thermal ALD and corresponding results were discussed in [42]. Following that, the TEM results for the films grown with plasma-assisted ALD approach will be presented and compared with those obtained for the films grown by thermal ALD and submitted to the same annealing treatment.

### 3.3.4. Microstructure of the films

Fig. 9 represents, as an example, a cross-sectional view of the film grown with Er(tmhd)₃ and annealed at $T_A$ = 900 °C. The film thickness is about 72 nm (Fig. 9,a) that supports the thickness value obtained by the ellipsometry spectra fitting (Fig. 3,b).

Dark field and bright field TEM images show that only the upper part of the film (over a thickness of about 53 nm) is crystalline. The dark field image was obtained from (200) spots of Er₂O₃. The bottom part has a uniform contrast that is characteristic for an amorphous material. It is seen that the grains are elongated following the growth direction with a dimension of about 53 nm for the largest ones and their average width is about 15 nm (Fig. 9,a and b).

Fig. 9,c and d are HRTEM images of enlarged regions indicated in Fig. 9,b by rectangular dashed boxes. The bottom region (about 19 nm thick) is clearly amorphous since no lattice fringes are visible in Fig. 9,c. By contrast, the enlarged top surface of a typical grain is crystalline and the FFT pattern, shown in inset, is consistent with the Er₂O₃ cubic phase. One can also notice the significant film roughness corresponding to the top of the elongated grains as observed by AFM method (Fig. 5). The comparison of these data with FTIR spectra allows assuming that this amorphous layer is an Er silicate layer caused by Si diffusion from the substrate. This has been confirmed by EDX measurements (see below).

Fig. 10 shows HRTEM images of the Er₂O₃ film grown with the plasma-assisted approach and annealed at 1100 °C. The upper part of the film keeps its cubic Er₂O₃ structure (Fig. 10,a). Indeed, the FFT in

inset is consistent with a [111] projection of a bixbyite structure (centered cubic structure) with lattice parameter *a* = 10.53 Å corresponding to the cubic Er₂O₃ structure.

Fig. 10,b clearly evidences the presence of a thicker silica layer (about 7.5 nm) compared to that (about 4.5 nm) observed upon 900 °C anneal treatment. On the top of this silica layer, a crystalline Er₂SiO₅ layer (about 20 nm thick) with elongated grains lying parallel to the substrate is found. The FFT in inset attests of the monoclinic structure of this phase with the following characteristics: *a* = 14.32 Å, *b* = 6.69 Å, *c* = 10.35 Å and *β* = 122.33°. These results are in agreement with the XRD data. The appearance of some pores (voids or cavities) located between the Er₂O₃ and Er₂SiO₅ layers is also detected which can be caused by desorption of residual chemical species from the film upon annealing as well as stresses relaxation in the films. However, this effect is less pronounced for the films grown with the plasma-assisted approach than with thermal ALD [42].

To confirm the silicate nature of the bottom part of the film, EDX mapping were performed. Fig. 11 is a panel of chemical maps (Er, Si and O elements) of the cross section of the film annealed at 1100 °C. The STEM HAADF image in Fig. 11,a clearly shows the three different regions of the film: the bottom 7.5 nm thick silica layer, the intermediate erbium silicate layer (about 25 nm thick) and the upper erbium oxide layer (about 40 nm thick). Note the small dark regions in the middle of the film in Fig. 11,a. These latter are the voids evoked previously.

The various contrasts of the chemical maps and more particularly those of Er and Si maps, confirm the previous statements: the Er map shows a brighter region at the upper part of the film confirming the higher content in Er in this region (erbium oxide film). The Si map clearly depicts the presence of Si at the bottom part of the film which is consistent with the silica and the erbium silicate layers in this region. A quantitative approach reveals that one have approximately twice the amount of Er compared to the Si content which is consistent with the Er₂SiO₅ phase. These findings support the conclusions made from FTIR spectra analysis (Figs. 7 and 8) as well as from ellipsometry data (Fig. 3). The evolution of film morphology and crystalline structure can also affect light emitting properties of the films. In particular, photoluminescence will be discussed in the next section.
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Comparison of FTIR spectra of as-deposited (AD) and annealed at different temperatures Er₂O₃ films grown with Er(CpMe)₃ by the thermal ALD and with Er(tmhd)₃ by the plasma assisted ALD. The spectra were recorded with normal and 60° incidence of excitation light and were shifted vertically in the figures for the clarity. Annealing temperatures are mentioned in the figures.</figcaption></figure>

for the films annealed at $T_A$ = 1100 °C only whereas such process became to be significant already for the films grown with Er(tmhd)₃ precursor at $T_A$ = 900 °C (Figs. 7 and 8).

An analysis of these samples using the TEM technique is necessary to get insight on the spatial localization of SiOₓ and/or Er silicate phases in the samples annealed at high $T_A$. Earlier we have observed cross-sections of the films grown with thermal ALD and corresponding results were discussed in [42]. Following that, the TEM results for the films grown with plasma-assisted ALD approach will be presented and compared with those obtained for the films grown by thermal ALD and submitted to the same annealing treatment.

### 3.3.4. Microstructure of the films

Fig. 9 represents, as an example, a cross-sectional view of the film grown with Er(tmhd)₃ and annealed at $T_A$ = 900 °C. The film thickness is about 72 nm (Fig. 9,a) that supports the thickness value obtained by the ellipsometry spectra fitting (Fig. 3,b).

Dark field and bright field TEM images show that only the upper part of the film (over a thickness of about 53 nm) is crystalline. The dark field image was obtained from (200) spots of Er₂O₃. The bottom part has a uniform contrast that is characteristic for an amorphous material. It is seen that the grains are elongated following the growth direction with a dimension of about 53 nm for the largest ones and their average width is about 15 nm (Fig. 9,a and b).

Fig. 9,c and d are HRTEM images of enlarged regions indicated in Fig. 9,b by rectangular dashed boxes. The bottom region (about 19 nm thick) is clearly amorphous since no lattice fringes are visible in Fig. 9,c. By contrast, the enlarged top surface of a typical grain is crystalline and the FFT pattern, shown in inset, is consistent with the Er₂O₃ cubic phase. One can also notice the significant film roughness corresponding to the top of the elongated grains as observed by AFM method (Fig. 5). The comparison of these data with FTIR spectra allows assuming that this amorphous layer is an Er silicate layer caused by Si diffusion from the substrate. This has been confirmed by EDX measurements (see below).

Fig. 10 shows HRTEM images of the Er₂O₃ film grown with the plasma-assisted approach and annealed at 1100 °C. The upper part of the film keeps its cubic Er₂O₃ structure (Fig. 10,a). Indeed, the FFT in

inset is consistent with a [111] projection of a bixbyite structure (centered cubic structure) with lattice parameter *a* = 10.53 Å corresponding to the cubic Er₂O₃ structure.

Fig. 10,b clearly evidences the presence of a thicker silica layer (about 7.5 nm) compared to that (about 4.5 nm) observed upon 900 °C anneal treatment. On the top of this silica layer, a crystalline Er₂SiO₅ layer (about 20 nm thick) with elongated grains lying parallel to the substrate is found. The FFT in inset attests of the monoclinic structure of this phase with the following characteristics: *a* = 14.32 Å, *b* = 6.69 Å, *c* = 10.35 Å and *β* = 122.33°. These results are in agreement with the XRD data. The appearance of some pores (voids or cavities) located between the Er₂O₃ and Er₂SiO₅ layers is also detected which can be caused by desorption of residual chemical species from the film upon annealing as well as stresses relaxation in the films. However, this effect is less pronounced for the films grown with the plasma-assisted approach than with thermal ALD [42].

To confirm the silicate nature of the bottom part of the film, EDX mapping were performed. Fig. 11 is a panel of chemical maps (Er, Si and O elements) of the cross section of the film annealed at 1100 °C. The STEM HAADF image in Fig. 11,a clearly shows the three different regions of the film: the bottom 7.5 nm thick silica layer, the intermediate erbium silicate layer (about 25 nm thick) and the upper erbium oxide layer (about 40 nm thick). Note the small dark regions in the middle of the film in Fig. 11,a. These latter are the voids evoked previously.

The various contrasts of the chemical maps and more particularly those of Er and Si maps, confirm the previous statements: the Er map shows a brighter region at the upper part of the film confirming the higher content in Er in this region (erbium oxide film). The Si map clearly depicts the presence of Si at the bottom part of the film which is consistent with the silica and the erbium silicate layers in this region. A quantitative approach reveals that one have approximately twice the amount of Er compared to the Si content which is consistent with the Er₂SiO₅ phase. These findings support the conclusions made from FTIR spectra analysis (Figs. 7 and 8) as well as from ellipsometry data (Fig. 3). The evolution of film morphology and crystalline structure can also affect light emitting properties of the films. In particular, photoluminescence will be discussed in the next section.
<figure><img src="image_10.png" /><figcaption>**Fig. 9.** Dark (a) and bright (b) fields TEM images of a cross sectional view of the Er₂O₃ film grown with the plasma-assisted approach and annealed at T_A = 900 °C. High-resolution TEM images of the bottom region of the film (c) and the top of the film (d). The inserted FFT pattern in (d) is in agreement with the cubic structure of Er₂O₃.</figcaption></figure>

### 3.3.5. Photoluminescence properties of the films

Light emission of bulk erbium oxide is caused by the native defects (for instance, oxygen vacancies) [40,41,42]. Besides, Er³⁺ ions can give the contribution in luminescence spectrum. In this regard, light emitting properties of the films investigated will be considered taking into account contribution of carrier recombination via defects and via Er³⁺ ions. The difference between them can be made using different excitation light wavelengths. Indeed, the use of non-resonant excitation for Er³⁺ ions can allow separating PL bands related to the Er₂O₃ host defects and Er³⁺ ions. Besides, it the use of non-resonant excitation for Er³⁺ ions can help to determine the role of the defects as the sensitizers of Er³⁺ ions similar to the case described in Refs. [39,46]. The PL emission of Er³⁺ ions can be directly studied under resonant excitation. The comparison of these approaches will permit us to get insight into the PL mechanism in the films investigated.

To observe Er³⁺ PL emission, the spectra were recorded under "resonant" excitation corresponding to the Er³⁺ absorption at 488 nm

<figure><img src="image_11.png" /><figcaption>**Fig. 10.** (a) HRTEM image taken from the upper part of the film grown with the plasma-assisted approach and annealed at 1100 °C. It shows the presence of crystalline grains having the cubic structure of Er₂O₃. (b) HRTEM image of the bottom region of the film revealing the presence of a crystalline Er silicate phase.</figcaption></figure>

(⁴I₁₅/₂→⁴F₇/₂ transition) and 532 nm (⁴I₁₅/₂→⁴H₁₁/₂ transition). Two main bands in the visible range were detected due to ⁴F₉/₂→⁴I₁₅/₂ (650–700 nm) and ⁴I₉/₂→⁴I₁₅/₂ (780–830 nm) optical transitions [48].

In the as-deposited films, the Er³⁺ related peaks were broad testifying some disorder in the film volume. This broadening was more pronounced in the films grown with thermal ALD approach [42]. Annealing at T_A = 500–900 °C show sharp Er³⁺ peaks (Fig. 12,a and b) that confirms the presence of Er³⁺ ions in the crystalline host. Further T_A rise up to 1100 °C results in the appearance of broad PL bands peaked near 620 and 700 nm, being overlapped with Er³⁺ PL bands (Fig. 12).

The comparison of the evolution of structural and light emitting properties of the films permitted to assume that these broad PL components are due to some defects of Er silicate phase which formation was more pronounced in the films grown with the plasma assisted ALD approach. This finding is supported by the FTIR data and structural evolution discussed above.

According to PL data, there are, at least, two different excitation and radiative channels in the films studied. One of them is ascribed to Er₂O₃
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Dark (a) and bright (b) fields TEM images of a cross sectional view of the Er₂O₃ film grown with the plasma-assisted approach and annealed at T_A = 900 °C. High-resolution TEM images of the bottom region of the film (c) and the top of the film (d). The inserted FFT pattern in (d) is in agreement with the cubic structure of Er₂O₃.</figcaption></figure>

### 3.3.5. Photoluminescence properties of the films

Light emission of bulk erbium oxide is caused by the native defects (for instance, oxygen vacancies) [40,41,42]. Besides, Er³⁺ ions can give the contribution in luminescence spectrum. In this regard, light emitting properties of the films investigated will be considered taking into account contribution of carrier recombination via defects and via Er³⁺ ions. The difference between them can be made using different excitation light wavelengths. Indeed, the use of non-resonant excitation for Er³⁺ ions can allow separating PL bands related to the Er₂O₃ host defects and Er³⁺ ions. Besides, it the use of non-resonant excitation for Er³⁺ ions can help to determine the role of the defects as the sensitizers of Er³⁺ ions similar to the case described in Refs. [39,46]. The PL emission of Er³⁺ ions can be directly studied under resonant excitation. The comparison of these approaches will permit us to get insight into the PL mechanism in the films investigated.

To observe Er³⁺ PL emission, the spectra were recorded under "resonant" excitation corresponding to the Er³⁺ absorption at 488 nm

<figure><img src="image_12.png" /><figcaption>**Fig. 10.** (a) HRTEM image taken from the upper part of the film grown with the plasma-assisted approach and annealed at 1100 °C. It shows the presence of crystalline grains having the cubic structure of Er₂O₃. (b) HRTEM image of the bottom region of the film revealing the presence of a crystalline Er silicate phase.</figcaption></figure>

(⁴I₁₅/₂→⁴F₇/₂ transition) and 532 nm (⁴I₁₅/₂→⁴H₁₁/₂ transition). Two main bands in the visible range were detected due to ⁴F₉/₂→⁴I₁₅/₂ (650–700 nm) and ⁴I₉/₂→⁴I₁₅/₂ (780–830 nm) optical transitions [48].

In the as-deposited films, the Er³⁺ related peaks were broad testifying some disorder in the film volume. This broadening was more pronounced in the films grown with thermal ALD approach [42]. Annealing at T_A = 500–900 °C show sharp Er³⁺ peaks (Fig. 12,a and b) that confirms the presence of Er³⁺ ions in the crystalline host. Further T_A rise up to 1100 °C results in the appearance of broad PL bands peaked near 620 and 700 nm, being overlapped with Er³⁺ PL bands (Fig. 12).

The comparison of the evolution of structural and light emitting properties of the films permitted to assume that these broad PL components are due to some defects of Er silicate phase which formation was more pronounced in the films grown with the plasma assisted ALD approach. This finding is supported by the FTIR data and structural evolution discussed above.

According to PL data, there are, at least, two different excitation and radiative channels in the films studied. One of them is ascribed to Er₂O₃
<figure><img src="image_12.png" /><figcaption>Fig. 11. (a) STEM HAADF image of the cross section of the Er₂O₃ film grown with the plasma-assisted approach and annealed at 1100 °C; (b) Er, (c) Si and (d) O chemical maps of the same region.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>Fig. 12. PL spectra of as-deposited and annealed films grown with plasma-assisted (a) and thermal (b) approaches. The **luminescence** was excited with a 532-nm line of Nd:YAG laser source at room temperature.</figcaption></figure>
<figure><img src="image_13.png" /><figcaption>Fig. 11. (a) STEM HAADF image of the cross section of the Er₂O₃ film grown with the plasma-assisted approach and annealed at 1100 °C; (b) Er, (c) Si and (d) O chemical maps of the same region.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>Fig. 12. PL spectra of as-deposited and annealed films grown with plasma-assisted (a) and thermal (b) approaches. The **luminescence** was excited with a 532-nm line of Nd:YAG laser source at room temperature.</figcaption></figure>
host (for instance, its native defects as oxygen vacancies) while the other
can be linked to intra-4f shell transitions of Er³⁺ ions [42]. The contri-
bution of Er₂O₃ host defects in PL spectrum is significant under ultra-
violet excitation. At the same time, the use of visible excitations (for
instance, 488 or 532 nm) permits the contribution of Er³⁺ ions in PL
spectrum to be revealed (Fig. 12).

Fig. 13 shows PL spectra recorded for Er³⁺ and more particularly its
⁴I₁₃/₂→⁴I₁₅/₂ optical transition where the contribution of oxygen va-
cancies was prevented. It is seen that with Tᴀ rise up to 1000 °C, Er³⁺ PL
intensity increases showing a saturation behavior. However, further Tᴀ
increase (up to 1100 °C) results in Er³⁺ PL enhancement. This effect can
be caused by the increase of Er³⁺ neighbor distance in Er silicate phase
and, as a consequence, by the reduction of the concentration quenching
of Er³⁺ PL emission. Another reason for PL enhancement could be the
rise in the number of emitting centers, i.e. the increase of the number of
“optically active” Er³⁺ ions in the silicate phase.

In order to obtain information about the contribution of host defects, PL spectra were recorded under ultraviolet excitation (Fig.14a). It is seen that under the excitation with 360 nm light (non-resonant for Er³⁺ ions), the PL spectrum is broad and featureless and it can be ascribed to the Er₂O₃ host defects as oxygen vacancies. At the same time, under resonant excitation (378 nm), corresponding to the ⁴I₁₅/₂→⁴G₁₁/₂ absorption for Er³⁺ ions, permit to detect sharp peaks related to Er³⁺ PL emission due to ²G₇/₂, ²G₉/₂→⁴I₁₅/₂, ²K₁₅/₂→⁴I₁₅/₂ and ²H₉/₂→⁴I₁₅/₂ transitions. Similar Er³⁺ peaks are detected under excitation with 488 and 520 nm (Figs. 12 and 13).

Additional arguments towards the potential energy transfers from
host defects to Er³⁺ ions can be obtained from PL excitation spectra
recorded for the same films (Fig. 14,b). It turns out that the PLE spectra
recorded for 550 and 564 nm detection wavelengths (Fig. 14,b, curves
1,2) are similar to the optical absorption of Er₂O₃ powder (Fig. 14,b,
curve 3) that is the evidence of the direct excitation via Er³⁺ ions.
Furthermore, some broad background in the PLE spectra is also distin-
guished that may originate from some contribution of oxygen vacancies
in the excitation process . The comparison of absorption spectrum and
PLE ones recorded at 400 nm wavelength is shown for the films annealed
at 1000°C (Fig.14,b, inset). The spectra's behaviors are similar, with the

exception that the PLE spectrum has some additional bands (or shoul-
ders) near 270 and 320 nm. These bands could be caused by self-trapped
excitons and host defects, such as oxygen vacancies, complexes of de-
fects, etc. that can contribute in PL emission via charge-transfer band.
These PLE features are similar to those observed for other Er-doped
oxides and silicates [21,42]. However, no significant energy transfer is
observed between vacancies and Er³⁺ ions meaning that native defects
of Er₂O₃ host were not efficient sensitizers of Er³⁺ emission in the films
investigated.Besides, in the films grown with thermal approach, the
presence of residual OH groups was observed in as-deposited films. It is
known that the Er³⁺ PL emission can be quenched via Er³⁺⁻OH inter-
action [49]. At high hydroxyl concentrations, direct resonant energy
transfer from the excited Er ion towards OH served as an effective traps.
At low OH concentration, fast energy transfer from ion to ion via the
Forster mechanism could cause the migration of the excitation towards
OH⁻ groups, where multi-phonon assisted decay to the oxide matrix.
Besides, rare-earth concentration is one of dominant factors for coop-
erative energy transfer. In our films grown with thermal approach both
possibilities could be realized because of high Er content. Since
annealing treatment results in the reduction of the number of hydroxyl
groups, one can expect increasing in PL intensity of Er³⁺ ions. However,
even small amount of OH groups at 900-1000°C (below 10⁻⁴ molar
fraction) can quench Er³⁺ PL emission significantly [49]. This fact could
explain higher intensity of rare-earth PL in the films grown with plasma
assisted approach. Apart from this, high-temperature annealing stimu-
lated formation of Er silicate phase. An appearance of weak oxygen
bonds, neutral oxygen vacancies, and Si=O states result in the increasing
of the PL emission in the 320-600 nm spectral range. Both PL and PLE
spectra of these defects are broad; as a consequence, these defects can be
efficient sensitizers of Er³⁺ PL emission (Fig. 13). This assumption is in
agreement with the results reported for the Er³⁺ emission in Er-doped
silicon rich oxide in Ref. [50].

4. Conclusions

This work demonstrates the influence of the nature of Er precursors
on the morphology and optical properties of Er₂O₃ films by atomic layer

<figure><img src="image_14.png" /><figcaption>**Fig. 13.** Effect of annealing temperature on the Er³⁺ PL intensity of the films grown with plasma-assisted ALD. PL evolution is shown for ⁴I₁₃/₂→⁴I₁₅/₂ transition. PL spectra of the films annealed at 500, 1000 and 1100 °C are presented in the inset. The spectra were recorded at room temperature with a 488 nm excitation light source.</figcaption></figure>
host (for instance, its native defects as oxygen vacancies) while the other
can be linked to intra-4f shell transitions of Er³⁺ ions [42]. The contri-
bution of Er₂O₃ host defects in PL spectrum is significant under ultra-
violet excitation. At the same time, the use of visible excitations (for
instance, 488 or 532 nm) permits the contribution of Er³⁺ ions in PL
spectrum to be revealed (Fig. 12).

Fig. 13 shows PL spectra recorded for Er³⁺ and more particularly its
⁴I₁₃/₂→⁴I₁₅/₂ optical transition where the contribution of oxygen va-
cancies was prevented. It is seen that with Tᴀ rise up to 1000 °C, Er³⁺ PL
intensity increases showing a saturation behavior. However, further Tᴀ
increase (up to 1100 °C) results in Er³⁺ PL enhancement. This effect can
be caused by the increase of Er³⁺ neighbor distance in Er silicate phase
and, as a consequence, by the reduction of the concentration quenching
of Er³⁺ PL emission. Another reason for PL enhancement could be the
rise in the number of emitting centers, i.e. the increase of the number of
“optically active” Er³⁺ ions in the silicate phase.

In order to obtain information about the contribution of host defects, PL spectra were recorded under ultraviolet excitation (Fig.14a). It is seen that under the excitation with 360 nm light (non-resonant for Er³⁺ ions), the PL spectrum is broad and featureless and it can be ascribed to the Er₂O₃ host defects as oxygen vacancies. At the same time, under resonant excitation (378 nm), corresponding to the ⁴I₁₅/₂→⁴G₁₁/₂ absorption for Er³⁺ ions, permit to detect sharp peaks related to Er³⁺ PL emission due to ²G₇/₂, ²G₉/₂→⁴I₁₅/₂, ²K₁₅/₂→⁴I₁₅/₂ and ²H₉/₂→⁴I₁₅/₂ transitions. Similar Er³⁺ peaks are detected under excitation with 488 and 520 nm (Figs. 12 and 13).

Additional arguments towards the potential energy transfers from
host defects to Er³⁺ ions can be obtained from PL excitation spectra
recorded for the same films (Fig. 14,b). It turns out that the PLE spectra
recorded for 550 and 564 nm detection wavelengths (Fig. 14,b, curves
1,2) are similar to the optical absorption of Er₂O₃ powder (Fig. 14,b,
curve 3) that is the evidence of the direct excitation via Er³⁺ ions.
Furthermore, some broad background in the PLE spectra is also distin-
guished that may originate from some contribution of oxygen vacancies
in the excitation process . The comparison of absorption spectrum and
PLE ones recorded at 400 nm wavelength is shown for the films annealed
at 1000°C (Fig.14,b, inset). The spectra's behaviors are similar, with the

exception that the PLE spectrum has some additional bands (or shoul-
ders) near 270 and 320 nm. These bands could be caused by self-trapped
excitons and host defects, such as oxygen vacancies, complexes of de-
fects, etc. that can contribute in PL emission via charge-transfer band.
These PLE features are similar to those observed for other Er-doped
oxides and silicates [21,42]. However, no significant energy transfer is
observed between vacancies and Er³⁺ ions meaning that native defects
of Er₂O₃ host were not efficient sensitizers of Er³⁺ emission in the films
investigated.Besides, in the films grown with thermal approach, the
presence of residual OH groups was observed in as-deposited films. It is
known that the Er³⁺ PL emission can be quenched via Er³⁺⁻OH inter-
action [49]. At high hydroxyl concentrations, direct resonant energy
transfer from the excited Er ion towards OH served as an effective traps.
At low OH concentration, fast energy transfer from ion to ion via the
Forster mechanism could cause the migration of the excitation towards
OH⁻ groups, where multi-phonon assisted decay to the oxide matrix.
Besides, rare-earth concentration is one of dominant factors for coop-
erative energy transfer. In our films grown with thermal approach both
possibilities could be realized because of high Er content. Since
annealing treatment results in the reduction of the number of hydroxyl
groups, one can expect increasing in PL intensity of Er³⁺ ions. However,
even small amount of OH groups at 900-1000°C (below 10⁻⁴ molar
fraction) can quench Er³⁺ PL emission significantly [49]. This fact could
explain higher intensity of rare-earth PL in the films grown with plasma
assisted approach. Apart from this, high-temperature annealing stimu-
lated formation of Er silicate phase. An appearance of weak oxygen
bonds, neutral oxygen vacancies, and Si=O states result in the increasing
of the PL emission in the 320-600 nm spectral range. Both PL and PLE
spectra of these defects are broad; as a consequence, these defects can be
efficient sensitizers of Er³⁺ PL emission (Fig. 13). This assumption is in
agreement with the results reported for the Er³⁺ emission in Er-doped
silicon rich oxide in Ref. [50].

4. Conclusions

This work demonstrates the influence of the nature of Er precursors
on the morphology and optical properties of Er₂O₃ films by atomic layer

<figure><img src="image_15.png" /><figcaption>**Fig. 13.** Effect of annealing temperature on the Er³⁺ PL intensity of the films grown with plasma-assisted ALD. PL evolution is shown for ⁴I₁₃/₂→⁴I₁₅/₂ transition. PL spectra of the films annealed at 500, 1000 and 1100 °C are presented in the inset. The spectra were recorded at room temperature with a 488 nm excitation light source.</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>**Fig. 14.** (a) PL spectra of Er₂O₃ film prepared with plasma-assisted approach and annealed at 1000 °C. The excitation wavelength is 360 nm (1), 378 nm (2), 488 nm (3) and 520 nm (4); (b) PLE spectra of the same film detected at 400, 550 and 564 nm. The spectrum of optical absorption of Er₂O₃ powder is given for comparison. Inset shows absorption spectra (1,2) and PLE spectrum detected at 400 nm wavelength (3) of the as-deposited film (1) and annealed at 1000 °C (2,3).</figcaption></figure>

deposition and submitted to high-temperature annealing. This latter stimulates pronounced film crystallization accompanied by the formation of Er silicate layer near film/substrate interface due to Si out-diffusion from underlying Si substrate. Upon an annealing at 1000–1100 °C, the appearance of Er₂SiO₅ phase in film volume occurs. Light emitting properties of the films are controlled by the competition of carrier recombination via Er₂O₃ host defects (like oxygen vacancies) as well as via intra-4f shell transitions in Er³⁺ ions. Formation of Er₂SiO₅ phase leads to an enhancement of Er³⁺ emission. Comparison of the evolution of Er³⁺ PL emission with annealing temperature for both types of Er₂O₃ films shows the presence of sharper Er³⁺ PL peaks in the films grown with plasma assisted approach. This is an additional arguments towards more prominent crystallization process in the films grown with Er(tmhd)₃ precursor.

### CRediT authorship contribution statement

**L. Khomenkova:** Conceptualization, Methodology, Investigation, Data curation, Writing – original draft, Writing – review & editing. **H. Merabet:** Funding acquisition, Supervision, Project administration, Validation, Writing – review & editing. **M.-P. Chauvat:** Data curation, Investigation, Visualization. **C. Frilay:** Resources. **X. Portier:** Data curation, Investigation, Visualization, Validation, Writing – review & editing. **C. Labbe:** Investigation, Validation, Writing – review & editing. **P. Marie:** Data curation, Formal analysis, Visualization. **J. Cardin:** Investigation, Formal analysis, Visualization, Validation, Writing – review & editing. **S. Boudin:** Resources. **J.-M. Rueff:** Data curation, Investigation, Visualization, Validation, Writing – review & editing. **F. Gourbilleau:** Funding acquisition, Supervision, Project administration, Validation, Writing – review & editing.

### Declaration of Competing Interest

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

### Acknowledgements

This work was partially supported by the Normandy Region and European Regional Development Fund (Fonds Européen de Développement Régional (FEDER)) via LUMIERE project as well as partially supported by the CNRS Federation IRMA - FR 3095 and partly supported by the Qatar National Research Fund (QNRF, a member of the Qatar Foundation) via NPRP grant 8-1467-1-268. The findings achieved herein are solely the responsibility of the authors. The authors would like to thank Dr. Nicolas Ratel-Ramond (CEMES, Toulouse, France) for the XRD experiments. LK and XP also thanks Franck Lemarié for the samples' preparation by FIB for TEM experiments. XP also acknowledges the "Agence Nationale de la Recherche" (ANR) for the EQUIPEX "GENESIS" grant "ANR-11-EQPX-0020" in the frame of the "Investissements d'avenir" whose funding allowed the acquisition of the FIB equipment.

### Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.surfin.2022.102377.

### References

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[2] L. Hu, W. Qi, Y. Li, Coatings strategies for atomic layer deposition, Nanotechnol. Rev. 6 (2017) 527-547.

[3] H.C.M. Knoops, T. Faraz, K. Arts, W.M.M. (Erwin) Kessels, Status and prospects of plasma-assisted atomic layer deposition, J. Vac. Sci. Technol. A 37 (2019), 030902.

[4] S.M. George, Atomic layer deposition: an overview, Chem. Rev. 110 (2010) 111-131.

[5] K.H. Goh, A.S.M.A. Haseeb, Y.H. Wong, Lanthanide rare earth oxide thin film as an alternative gate oxide, Mater. Sci. Semicond. Process. 68 (2017) 302-315, https://doi.org/10.1016/j.mssp.2017.06.037.

[6] B.P. Gila, J.W. Johnson, R. Mehandru, B. Luo, A.H. Onstine, K.K. Allums, V. Krishnamoorthy, S. Bates, C.R. Abernathy, F. Ren, S.J. Pearton, Gadolinium oxide and scandium oxide: gate dielectrics for GaN MOSFETs, Phys. Status Solidi A 188 (2001) 239.
<figure><img src="image_16.png" /><figcaption>**Fig. 14.** (a) PL spectra of Er₂O₃ film prepared with plasma-assisted approach and annealed at 1000 °C. The excitation wavelength is 360 nm (1), 378 nm (2), 488 nm (3) and 520 nm (4); (b) PLE spectra of the same film detected at 400, 550 and 564 nm. The spectrum of optical absorption of Er₂O₃ powder is given for comparison. Inset shows absorption spectra (1,2) and PLE spectrum detected at 400 nm wavelength (3) of the as-deposited film (1) and annealed at 1000 °C (2,3).</figcaption></figure>

deposition and submitted to high-temperature annealing. This latter stimulates pronounced film crystallization accompanied by the formation of Er silicate layer near film/substrate interface due to Si out-diffusion from underlying Si substrate. Upon an annealing at 1000–1100 °C, the appearance of Er₂SiO₅ phase in film volume occurs. Light emitting properties of the films are controlled by the competition of carrier recombination via Er₂O₃ host defects (like oxygen vacancies) as well as via intra-4f shell transitions in Er³⁺ ions. Formation of Er₂SiO₅ phase leads to an enhancement of Er³⁺ emission. Comparison of the evolution of Er³⁺ PL emission with annealing temperature for both types of Er₂O₃ films shows the presence of sharper Er³⁺ PL peaks in the films grown with plasma assisted approach. This is an additional arguments towards more prominent crystallization process in the films grown with Er(tmhd)₃ precursor.

### CRediT authorship contribution statement

**L. Khomenkova:** Conceptualization, Methodology, Investigation, Data curation, Writing – original draft, Writing – review & editing. **H. Merabet:** Funding acquisition, Supervision, Project administration, Validation, Writing – review & editing. **M.-P. Chauvat:** Data curation, Investigation, Visualization. **C. Frilay:** Resources. **X. Portier:** Data curation, Investigation, Visualization, Validation, Writing – review & editing. **C. Labbe:** Investigation, Validation, Writing – review & editing. **P. Marie:** Data curation, Formal analysis, Visualization. **J. Cardin:** Investigation, Formal analysis, Visualization, Validation, Writing – review & editing. **S. Boudin:** Resources. **J.-M. Rueff:** Data curation, Investigation, Visualization, Validation, Writing – review & editing. **F. Gourbilleau:** Funding acquisition, Supervision, Project administration, Validation, Writing – review & editing.

### Declaration of Competing Interest

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

### Acknowledgements

This work was partially supported by the Normandy Region and European Regional Development Fund (Fonds Européen de Développement Régional (FEDER)) via LUMIERE project as well as partially supported by the CNRS Federation IRMA - FR 3095 and partly supported by the Qatar National Research Fund (QNRF, a member of the Qatar Foundation) via NPRP grant 8-1467-1-268. The findings achieved herein are solely the responsibility of the authors. The authors would like to thank Dr. Nicolas Ratel-Ramond (CEMES, Toulouse, France) for the XRD experiments. LK and XP also thanks Franck Lemarié for the samples' preparation by FIB for TEM experiments. XP also acknowledges the "Agence Nationale de la Recherche" (ANR) for the EQUIPEX "GENESIS" grant "ANR-11-EQPX-0020" in the frame of the "Investissements d'avenir" whose funding allowed the acquisition of the FIB equipment.

### Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.surfin.2022.102377.

### References

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[8] W.-H. Kim, W.J. Maeng, K.-J. Moon, H. Kim, Growth characteristics and electrical properties of La₂O₃ gate oxides grown by thermal and plasma-enhanced atomic layer deposition, Thin Solid Films 519 (2010) 362–366, https://doi.org/10.1016/j.tsf.2010.07.108.

[9] H.D.B. Gottlob, M.C. Lemme, T. Mollenhauer, T. Wahlbrink, J.K. Efavi, H. Kurz, Y. Stefanov, K. Haberle, R. Komaragiri, T. Ruland, F. Zaunert, U. Schwalke, Introduction of crystalline high-k gate dielectrics in a CMOS process, J. Non Cryst. Solids 351 (2005) 1885.

[10] W.C. Chin, K.Y. Cheong, Z. Hassan, Sm₂O₃ gate dielectric on Si substrate, Mater. Sci. Semicond. Process. 13 (2010) 303.

[11] M. Losurdo, M.M. Giangregorio, G. Bruno, D. Yang, E.A. Irene, A.A. Suvorova, M. Saunders, Er₂O₃ as a high-k dielectric candidate, Appl. Phys. Lett. 91 (2007), 091914.

[12] J. Päiväsaari, M. Putkonen, T. Sajavaara, L. Niinistö, Atomic layer deposition of rare earth oxides: erbium oxide thin films from β-diketonate and ozone precursors, J. Alloy. Compd. 374 (2004) 124–128.

[13] P.-Y. Chen, A.B. Posadas, S. Kwon, Q. Wang, M.J. Kim, A.A. Demkov, J.G. Ekerdt, Cubic crystalline erbium oxide growth on GaN(0001) by atomic layer deposition, J. Appl. Phys. 122 (2017), 215302.

[14] S. Ohmi, C. Kobayashi, I. Kashiwagi, C. Ohshima, H. Ishiwara, H. Iwai, Characterization of La₂O₃ and Yb₂O₃ thin films for high-k gate insulator application, J. Electrochem. Soc. 150 (2003) F134–F140.

[15] K.M. Hubbard, B.F. Espinoza, Corrosion-resistant erbium oxide coatings by organometallic chemical vapor deposition, Thin Solid Films 366 (2000) 175–180.

[16] V. Mikhelashvili, G. Eisenstein, F. Edelmann, Characteristics of electron-beam-gun-evaporated Er₂O₃ thin films as gate dielectrics for silicon, J. Appl. Phys. 90 (2001) 5447–5449.

[17] D. Xue, K. Betzler, H. Hesse, Dielectric constants of binary rare-earth compounds, J. Phys. Condens. Matter 12 (2000) 3113–3118.

[18] S. Saini, K. Chen, X. Duan, J. Michel, L.C. Kimerling, M. Lipson, Er₂O₃ for high-gain waveguide amplifiers, J. Electron. Mater. 33 (2004) 809–814.

[19] M. Miritello, R. Lo Savio, A.M. Piro, G. Franzò, F. Priolo, F. Iacona, C. Bongiorno, Optical and structural properties of Er₂O₃ films grown by magnetron sputtering, J. Appl. Phys. 100 (2006), 013502.

[20] T.M. Pan, C.L. Chen, W.W. Yeh, S.J. Hou, Structural and electrical characteristics of thin erbium oxide gate dielectrics, Appl. Phys. Lett. 89 (2006), 222912.

[21] M. Miritello, R. Lo Savio, F. Iacona, G. Franzò, C. Bongiorno, A. Irrera, F. Priolo, The influence of substrate on the properties of Er₂O₃ films grown by magnetron sputtering, J. Lumin. 121 (2006) 233–237.

[22] M.P. Singh, C.S. Thakur, K. Shalini, N. Bhat, S.A. Shivashankar, Structural and electrical characterization of erbium oxide films grown on Si (100) by low-pressure metalorganic chemical vapor deposition, Appl. Phys. Lett. 83 (2003) 2889–2891.

[23] Y.H. Chang, M.H. Chou, T.J. Wang, Synthesis and photoluminescence properties of erbium oxide thin films prepared by sol-gel method, Ceram. Int. 44 (2018) 163–1167.

[24] M. Miritello, R. Lo Savio, F. Iacona, G. Franzo, A. Irrera, A.M. Piro, C. Bongiorno, F. Priolo, Efficient luminescence and energy transfer in erbium silicate thin films, Adv. Mater. 19 (2007) 1582–1588.

[25] T.H. Phung, D.K. Srinivasan, P. Steinmann, R. Wise, M.B. Yu, Y.C. Yeo, C. Zhu, High performance metal-insulator-metal capacitors with Er₂O₃ on ALD SiO₂ for RF applications, J. Electrochem. Soc. 158 (2011) H1289–H1292.

[26] K. Xu, A.R. Chaudhuri, H. Parala, D. Schwendt, T. de los Arcos, H. Jörg Ostenb, A. Devi, Atomic layer deposition of Er₂O₃ thin films from Er tris-guanidinate and water: process optimization, film analysis and electrical properties, J. Mater. Chem. C 1 (2013) 3939–3946.

[27] STREM (Europe), France; https://www.strem.com/; for Er(CpMe₃)₃ see https://www.strem.com/catalog/v/68-8740/19/erbium_39470-10-5; for Er (tmdh)₃ see https://www.strem.com/catalog/v/68-8750/erbium_35733-23-4, 2022.

[28] K.J. Eisentraut, R.E. Sievers, Thermogravimetric studies of metal β-diketonates, J. Inorg. Nucl. Chem. 29 (1967) 1931–1936, https://doi.org/10.1016/0022-1902 (67)80452-4.

[29] The Crystallography Open Database (COD): an open-access collection of crystal structures, http://www.crystallography.net/cod/COD; Reference code for XRD data of Er₂O₃ is 96-154-1746; compound name: 1541745; 2022.

[30] Joint Committee for Powder Diffraction Data; PDF Card # 01-074-1983.

[31] V.P. Rao, B. Besancon, V. Omarjee, C. Dussarrat, Development of lanthanide precursors as dopants for advanced high-k materials, ECS Trans. 33 (2010) 145–156.

[32] J. Niinistö, T. Hatanpää, M. Kariniemi, M. Mäntymäki, L. Costelle, K. Mizohata, K. Kukli, M. Ritala, M. Leskelä, Cycloheptatrienyl-cyclopentadienyl heteroleptic precursors for atomic layer deposition of group 4 oxide thin films, Chem. Mater. 24 (2012) 2002–2008.

[33] S. Seppälä, M. Vehkamäki, K. Mizohata, W. Noh, J. Räisänen, M. Ritala, M. Leskelä, Comparative study on the use of novel heteroleptic cyclopentadienyl-based zirconium precursors with H₂O and O₃ for atomic layer deposition of ZrO₂, J. Vac. Sci. Technol. A 37 (2019), 010912.

[34] In this work, the R-200-Advanced ALD setup from Picosun Oy (Finland) was used. More details about PICOSUN R-series atomic layer deposition reactors can be found at https://www.picosun.com/product/r-200-advanced/, 2022 [Online].

[35] In this work, silicon wafers from "Sil'tronix Silicon Technologies" (France) were used. More details available at https://www.sil-tronix-st.com/en/silicon-wafer/sili-con-wafer-specifications, 2022.

[36] L. Khomenkova, X. Portier, J. Cardin, F. Gourbilleau, Thermal stability of high-k Si-rich Hf₂O₃ layers grown by RF magnetron sputtering, Nanotechnology 21 (2010), 285707 (12pp).

[37] L. Khomenkova, C. Labbe, X. Portier, M. Carrada, F. Gourbilleau, Undoped and Nd³⁺ doped Si-based single layers and superlattices for photonic applications, Phys. Status Solidi A 210 (2013) 1532–1543, https://doi.org/10.1002/pssa.201200942.

[38] L. Khomenkova, N. Korsunska, C. Labbe, X. Portier, F. Gourbilleau, The peculiarities of structural and optical properties of Hf₂O₃-based films co-doped with silicon and erbium, Appl. Surf. Sci. 471 (2019) 521–527.

[39] M.F. Al-Kuhaili, S.M.A. Durrani, Optical properties of erbium oxide thin films deposited by electron beam evaporation, Thin Solid Films 515 (2007) 2885–2890.

[40] H.S. Kamineni, V.K. Kamineni, R.L. Moore II, S. Gallis, A.C. Diebold, M. Huang, A. E. Kaloyeros, Optical and structural characterization of thermal oxidation effects of erbium thin films deposited by electron beam on silicon, J. Appl. Phys. 111 (2012), 013104.

[41] G. Adachi, N. Imanaka, The binary rare earth oxides, Chem. Rev. 98 (1998) 1479–1514.

[42] Y.H. Chang, M.H. Chou, T.J. Wang, Synthesis and photoluminescence properties of erbium oxide thin films prepared by sol-gel method, Ceram. Int. 44 (2018) 1163–1167.

[43] L. Khomenkova, M.P. Chauvat, P. Marie, C. Frilay, F. Lemarié, S. Boudin, X. Portier, N. Ratel-Ramond, C. Labbé, J. Cardin, F. Gourbilleau, Thermally induced evolution of optical and structural properties of Er₂O₃ films grown on Si substrates by thermal atomic layer deposition, Mater. Lett. 263 (2020), 127216.

[44] H. Ono, T. Katsumata, Interfacial reactions between thin rare-earth-metal oxide films and Si substrates, Appl. Phys. Lett. 78 (2001) 1832–1834.

[45] H. Wong, H. Iwai, On the scaling issues and high-k replacement of ultrathin gate dielectrics for nanoscale MOS transistors, Microelectron. Eng. 83 (2006) 1867–1904.

[46] W. Deqi, Y. Jincheng, Z. Hongsheng, C. Aimin, Li Feng, Leakage current mechanisms of ultrathin high-k Er₂O₃ gate dielectric film, J. Semicond. 30 (2009), 103003.

[47] J. Wang, Y. Xia, Y. Shi, Z. Shi, L. Pu, R. Zhang, Y. Zheng, Z. Tao, F. Lu, 1.54 µm photoluminescence emission and oxygen vacancy as sensitizer in Er-doped Hf₂O₃ films, Appl. Phys. Lett. 91 (2007), 191115.

[48] M. Dammak, D.L. Zhang, Spectra and energy levels of Er³⁺ in Er₂O₃ powder, J. Alloy. Compd. 407 (2006) 8–15.

[49] S. Abedrabbo, C. Lahlouh, A.T. Fiory, Analytical study of thermal annealing behaviour of erbium emission in Er₂O₃-sol-gel silica films, J. Phys. D Appl. Phys. 44 (2011) 315401.

[50] L. Jin, D. Li, L. Xiang, F. Wang, D. Yang, D. Que, Energy transfer from luminescent centers to Er³⁺ in erbium-doped silicon-rich oxide films, Nanoscale Res. Lett. 8 (2013) 366.
[7] C. Liu, E.F. Chor, L.S. Tan, A. Du, Epitaxial growth of Sc₂O₃ films on GaN (0001) by pulsed laser deposition, J. Vac. Sci. Technol. B 25 (2007) 754.

[8] W.-H. Kim, W.J. Maeng, K.-J. Moon, H. Kim, Growth characteristics and electrical properties of La₂O₃ gate oxides grown by thermal and plasma-enhanced atomic layer deposition, Thin Solid Films 519 (2010) 362–366, https://doi.org/10.1016/j.tsf.2010.07.108.

[9] H.D.B. Gottlob, M.C. Lemme, T. Mollenhauer, T. Wahlbrink, J.K. Efavi, H. Kurz, Y. Stefanov, K. Haberle, R. Komaragiri, T. Ruland, F. Zaunert, U. Schwalke, Introduction of crystalline high-k gate dielectrics in a CMOS process, J. Non Cryst. Solids 351 (2005) 1885.

[10] W.C. Chin, K.Y. Cheong, Z. Hassan, Sm₂O₃ gate dielectric on Si substrate, Mater. Sci. Semicond. Process. 13 (2010) 303.

[11] M. Losurdo, M.M. Giangregorio, G. Bruno, D. Yang, E.A. Irene, A.A. Suvorova, M. Saunders, Er₂O₃ as a high-k dielectric candidate, Appl. Phys. Lett. 91 (2007), 091914.

[12] J. Päiväsaari, M. Putkonen, T. Sajavaara, L. Niinistö, Atomic layer deposition of rare earth oxides: erbium oxide thin films from β-diketonate and ozone precursors, J. Alloy. Compd. 374 (2004) 124–128.

[13] P.-Y. Chen, A.B. Posadas, S. Kwon, Q. Wang, M.J. Kim, A.A. Demkov, J.G. Ekerdt, Cubic crystalline erbium oxide growth on GaN(0001) by atomic layer deposition, J. Appl. Phys. 122 (2017), 215302.

[14] S. Ohmi, C. Kobayashi, I. Kashiwagi, C. Ohshima, H. Ishiwara, H. Iwai, Characterization of La₂O₃ and Yb₂O₃ thin films for high-k gate insulator application, J. Electrochem. Soc. 150 (2003) F134–F140.

[15] K.M. Hubbard, B.F. Espinoza, Corrosion-resistant erbium oxide coatings by organometallic chemical vapor deposition, Thin Solid Films 366 (2000) 175–180.

[16] V. Mikhelashvili, G. Eisenstein, F. Edelmann, Characteristics of electron-beam-gun-evaporated Er₂O₃ thin films as gate dielectrics for silicon, J. Appl. Phys. 90 (2001) 5447–5449.

[17] D. Xue, K. Betzler, H. Hesse, Dielectric constants of binary rare-earth compounds, J. Phys. Condens. Matter 12 (2000) 3113–3118.

[18] S. Saini, K. Chen, X. Duan, J. Michel, L.C. Kimerling, M. Lipson, Er₂O₃ for high-gain waveguide amplifiers, J. Electron. Mater. 33 (2004) 809–814.

[19] M. Miritello, R. Lo Savio, A.M. Piro, G. Franzò, F. Priolo, F. Iacona, C. Bongiorno, Optical and structural properties of Er₂O₃ films grown by magnetron sputtering, J. Appl. Phys. 100 (2006), 013502.

[20] T.M. Pan, C.L. Chen, W.W. Yeh, S.J. Hou, Structural and electrical characteristics of thin erbium oxide gate dielectrics, Appl. Phys. Lett. 89 (2006), 222912.

[21] M. Miritello, R. Lo Savio, F. Iacona, G. Franzò, C. Bongiorno, A. Irrera, F. Priolo, The influence of substrate on the properties of Er₂O₃ films grown by magnetron sputtering, J. Lumin. 121 (2006) 233–237.

[22] M.P. Singh, C.S. Thakur, K. Shalini, N. Bhat, S.A. Shivashankar, Structural and electrical characterization of erbium oxide films grown on Si (100) by low-pressure metalorganic chemical vapor deposition, Appl. Phys. Lett. 83 (2003) 2889–2891.

[23] Y.H. Chang, M.H. Chou, T.J. Wang, Synthesis and photoluminescence properties of erbium oxide thin films prepared by sol-gel method, Ceram. Int. 44 (2018) 163–1167.

[24] M. Miritello, R. Lo Savio, F. Iacona, G. Franzo, A. Irrera, A.M. Piro, C. Bongiorno, F. Priolo, Efficient luminescence and energy transfer in erbium silicate thin films, Adv. Mater. 19 (2007) 1582–1588.

[25] T.H. Phung, D.K. Srinivasan, P. Steinmann, R. Wise, M.B. Yu, Y.C. Yeo, C. Zhu, High performance metal-insulator-metal capacitors with Er₂O₃ on AlD SiO₂ for RF applications, J. Electrochem. Soc. 158 (2011) H1289–H1292.

[26] K. Xu, A.R. Chaudhuri, H. Parala, D. Schwendt, T. de los Arcos, H. Jörg Ostenb, A. Devi, Atomic layer deposition of Er₂O₃ thin films from Er tris-guanidinate and water: process optimization, film analysis and electrical properties, J. Mater. Chem. C 1 (2013) 3939–3946.

[27] STREM (Europe), France; https://www.strem.com/; for Er(CpMe₃)₃ see https://www.strem.com/catalog/v/68-8740/19/erbium_39470-10-5; for Er (tmdh)₃ see https://www.strem.com/catalog/v/68-8750/erbium_35733-23-4, 2022.

[28] K.J. Eisentraut, R.E. Sievers, Thermogravimetric studies of metal β-diketonates, J. Inorg. Nucl. Chem. 29 (1967) 1931–1936, https://doi.org/10.1016/0022-1902 (67)80452-4.

[29] The Crystallography Open Database (COD): an open-access collection of crystal structures, http://www.crystallography.net/cod/COD; Reference code for XRD data of Er₂O₃ is 96-154-1746; compound name: 1541745; 2022.

[30] Joint Committee for Powder Diffraction Data; PDF Card # 01-074-1983.

[31] V.P. Rao, B. Besancon, V. Omarjee, C. Dussarrat, Development of lanthanide precursors as dopants for advanced high-k materials, ECS Trans. 33 (2010) 145–156.

[32] J. Niinistö, T. Hatanpää, M. Kariniemi, M. Mäntymäki, L. Costelle, K. Mizohata, K. Kukli, M. Ritala, M. Leskelä, Cycloheptatrienyl-cyclopentadienyl heteroleptic precursors for atomic layer deposition of group 4 oxide thin films, Chem. Mater. 24 (2012) 2002–2008.

[33] S. Seppälä, M. Vehkamäki, K. Mizohata, W. Noh, J. Räisänen, M. Ritala, M. Leskelä, Comparative study on the use of novel heteroleptic cyclopentadienyl-based zirconium precursors with H₂O and O₃ for atomic layer deposition of ZrO₂, J. Vac. Sci. Technol. A 37 (2019), 010912.

[34] In this work, the R-200-Advanced ALD setup from Picosun Oy (Finland) was used. More details about PICOSUN R-series atomic layer deposition reactors can be found at https://www.picosun.com/product/r-200-advanced/, 2022 [Online].

[35] In this work, silicon wafers from "Sil'tronix Silicon Technologies" (France) were used. More details available at https://www.sil-tronix-st.com/en/silicon-wafer/sili con-wafer-specifications, 2022.

[36] L. Khomenkova, X. Portier, J. Cardin, F. Gourbilleau, Thermal stability of high-k Si-rich Hf₂O₃ layers grown by RF magnetron sputtering, Nanotechnology 21 (2010), 285707 (12pp).

[37] L. Khomenkova, C. Labbe, X. Portier, M. Carrada, F. Gourbilleau, Undoped and Nd³⁺ doped Si-based single layers and superlattices for photonic applications, Phys. Status Solidi A 210 (2013) 1532–1543, https://doi.org/10.1002/pssa.201200942.

[38] L. Khomenkova, N. Korsunska, C. Labbe, X. Portier, F. Gourbilleau, The peculiarities of structural and optical properties of Hf₂O₃-based films co-doped with silicon and erbium, Appl. Surf. Sci. 471 (2019) 521–527.

[39] M.F. Al-Kuhaili, S.M.A. Durrani, Optical properties of erbium oxide thin films deposited by electron beam evaporation, Thin Solid Films 515 (2007) 2885–2890.

[40] H.S. Kamineni, V.K. Kamineni, R.L. Moore II, S. Gallis, A.C. Diebold, M. Huang, A. E. Kaloyeros, Optical and structural characterization of thermal oxidation effects of erbium thin films deposited by electron beam on silicon, J. Appl. Phys. 111 (2012), 013104.

[41] G. Adachi, N. Imanaka, The binary rare earth oxides, Chem. Rev. 98 (1998) 1479–1514.

[42] Y.H. Chang, M.H. Chou, T.J. Wang, Synthesis and photoluminescence properties of erbium oxide thin films prepared by sol-gel method, Ceram. Int. 44 (2018) 1163–1167.

[43] L. Khomenkova, M.P. Chauvat, P. Marie, C. Frilay, F. Lemarié, S. Boudin, X. Portier, N. Ratel-Ramond, C. Labbé, J. Cardin, F. Gourbilleau, Thermally induced evolution of optical and structural properties of Er₂O₃ films grown on Si substrates by thermal atomic layer deposition, Mater. Lett. 263 (2020), 127216.

[44] H. Ono, T. Katsumata, Interfacial reactions between thin rare-earth-metal oxide films and Si substrates, Appl. Phys. Lett. 78 (2001) 1832–1834.

[45] H. Wong, H. Iwai, On the scaling issues and high-k replacement of ultrathin gate dielectrics for nanoscale MOS transistors, Microelectron. Eng. 83 (2006) 1867–1904.

[46] W. Deqi, Y. Jincheng, Z. Hongsheng, C. Aimin, Li Feng, Leakage current mechanisms of ultrathin high-k Er₂O₃ gate dielectric film, J. Semicond. 30 (2009), 103003.

[47] J. Wang, Y. Xia, Y. Shi, Z. Shi, L. Pu, R. Zhang, Y. Zheng, Z. Tao, F. Lu, 1.54 µm photoluminescence emission and oxygen vacancy as sensitizer in Er-doped Hf₂O₃ films, Appl. Phys. Lett. 91 (2007), 191115.

[48] M. Dammak, D.L. Zhang, Spectra and energy levels of Er³⁺ in Er₂O₃ powder, J. Alloy. Compd. 407 (2006) 8–15.

[49] S. Abedrabbo, C. Lahlouh, A.T. Fiory, Analytical study of thermal annealing behaviour of erbium emission in Er₂O₃-sol-gel silica films, J. Phys. D Appl. Phys. 44 (2011) 315401.

[50] L. Jin, D. Li, L. Xiang, F. Wang, D. Yang, D. Que, Energy transfer from luminescent centers to Er³⁺ in erbium-doped silicon-rich oxide films, Nanoscale Res. Lett. 8 (2013) 366.