Oxygen vacancies dependent phase transition of Y2O3 films

Accepted Manuscript Title: Oxygen vacancies dependent phase transition of Y2 O3 films Author: Yu Pengfei Zhang Kan Huang Hao Wen Mao Li Quan Zhang Wei Hu Chaoquan Zheng Weitao PII: DOI: Reference:

S0169-4332(17)30823-1 http://dx.doi.org/doi:10.1016/j.apsusc.2017.03.145 APSUSC 35521

To appear in:

APSUSC

Received date: Revised date: Accepted date:

31-8-2016 13-2-2017 16-3-2017

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Highlights

Oxygen vacancies for Y2O3 films increase monotonously with increasing Ts.

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Oxygen vacancies can promote the nucleation of monoclinic phase. That monoclinic phase with oxygen deficiency is not thermodynamic stable at

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high temperature.

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Phase transition from monoclinic to oxygen defective occurs at high concentrations of oxygen vacancies.

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M

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High hardness just appears in Y2O3 films with mixed phase configurations.

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Oxygen vacancies dependent phase transition of Y2O3 films Yu Pengfei1, Zhang Kan1, Huang Hao2, Wen Mao1*, Li Quan1, Zhang Wei1, Hu Chaoquan1, Zheng Weitao3*

Department of Materials Science, State Key Laboratory of Superhard Materials, and

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1

People's Republic of China.

Titanium alloys lab. Beijing Institute of Aeronautical Materials, Beijing81-15 100095,

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2

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Key Laboratory of Automobile Materials, MOE, Jilin University, Changchun 130012,

3

an

People's Republic of China

Department of Materials Science, State Key Laboratory of Automotive Simulation

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and Control and Key Laboratory of Automobile Materials, MOE, Jilin University,

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Changchun 130012, People's Republic of China.

Abstract

Y2O3 films have great application potential in high-temperature metal matrix composite and nuclear engineering, used as interface diffusion and reaction barrier coating owing to their excellent thermal and chemical stability, high melting point and extremely negative Gibbs formation energy, and thus their structural and mechanical properties at elevated temperature are especially important. Oxygen vacancies exist commonly in yttrium oxide (Y2O3) thin films and act strongly on the phase structure and properties, but oxygen vacancies dependent phase transition at elevated temperature has not been well explored yet. Y2O3 thin films with different oxygen vacancy concentrations have been achieved by reactive sputtering through varying substrate temperature (Ts), in which oxygen vacancies increase monotonously with

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increasing Ts. For as-deposited Y2O3 films, oxygen vacancies present at high Ts can promote the nucleation of monoclinic phase, meanwhile, high Ts can induce the instability of monoclinic phase. Thus their competition results in forming mixed phases of cubic and monoclinic at high Ts. During vacuum annealing at 1000 °C, a

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critical oxygen vacancy concentration is observed, below which phase transition from monoclinic to cubic takes place, and above which phase transfer from monoclinic to

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from compressive to tensile and maintenance of high hardness.

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the oxygen defective phase (ICDD file no. 39-1063), accompanying by stress reversal

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Keywords: Y2O3 films; Oxygen vacancies; phase transition; hardness

*

Corresponding author, Tel./Fax: 86 431 85168246.

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Weitao).

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E-mail address: [email protected] (Wen Mao), [email protected] (Zheng

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1. Introduction Yttrium oxide (Y2O3) films have been exploited widely in various optical, electrical and electro-optic devices owing to their excellent electronic and optical properties,

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including high dielectric constant, superior electrical breakdown strength, high refractive index, low absorption, large band gap and a broad transmittance range[1-4].

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Moreover, excellent thermal and chemical stability[5] coupled with good mechanical

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properties[6] (high hardness and fracture toughness), high melting point and extremely negative Gibbs formation energy also make Y2O3 a suitable material for

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solid oxide fuel cells, nuclear engineering[2, 7-10], high-temperature protective

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coating[11], interface diffusion and reaction barrier coating of SiC fiber reinforced metal alloy matrix composite used as high-temperature structural components[12, 13].

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When Y2O3 films are applied in interface diffusion and reaction barrier coating

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between Si-based structural ceramics and alloys at elevated temperature, the temperature dependent structural and mechanical properties are of utmost importance, determining the interface reaction, adhesion and interfacial fracture. Y2O3 exhibits three typical structural polymorphisms: cubic, monoclinic, and

hexagonal, commonly known as C-, B-, and A-type structures, respectively[5]. The various properties of Y2O3 are strongly affected by its phase structure. For instance, B-type monoclinic phase at ordinary pressure owns lower density than C-type cubic phase[14, 15], and the photoluminescence properties of B-type monoclinic Y2O3 are quite different from those of C-type cubic phase[16-18]. To obtain tailored properties of Y2O3 for applications as interface diffusion and reaction barrier coating, it is crucial

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to control the phase structure and understand the structural stability at elevated temperature. For stoichiometric bulk Y2O3, its structural stability under high pressure and

ip t

temperature has been elucidated, in which the C-type cubic phase transforms to A-type hexagonal phase at about 2600 K or to A-type hexagonal phase through a

cr

two-step phase transition (cubic to monoclinic and monoclinic to hexagonal) by

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increasing the pressure up to 32.4 Gpa[5, 19, 20]. On the other hand, Y2O3 films have been prepared by a wide variety of technologies, including radio frequency

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sputtering[21-23], DC reactive magnetron sputtering[2], electron beam evaporation

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deposition[24], pulsed laser deposition[25, 26], ion-beam assisted deposition[11], and chemical vapor deposition[27, 28], in which nonstoichiometry related to crystal

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defects mostly appears in Y2O3 films due to far-from-equilibrium state and

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low-energy ions irradiation. The nature of defects makes it more difficult to control the phase structure in Y2O3 film by adjusting deposition parameters, and to elucidate structural stability by post annealing treatment. Substrate temperature (Ts) has been most widely applied to control the structure of Y2O3 films[22, 23, 25, 29-33], however, certain disagreements exist in different reports on the variation of phase structure and composition with Ts. As Ts is increased, some researchers[34] have observed the increment in the O/Y ratio and phase transition from B-type monoclinic to C-type cubic, whereas others[24, 31] have reported that O/Y ratio decreases and the phase maintains a cubic structure or transfers from monoclinic to cubic. Moreover, post annealing treatment can induce the phase transition from B-type monoclinic to

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C-type cubic, but the transition temperature varied from 400˚C to 1000˚C[11, 22, 25, 35-37]. Recently, it is reported that oxygen vacancies caused by ion bombardment can promote nucleation of B-type monoclinic phase both during[23] and after film

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growth[38] due to a local accommodation mechanism of the oxygen vacancies. In a word, the composition and defects play critical roles in controlling phase structure and

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determining the structural stability in Y2O3 films, which have not yet been well

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explored. The study of their oxygen vacancies dependent phase transition during deposition and after post annealing treatment is therefore of great importance and is

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the subject of the current paper, since oxygen vacancies have been considered as the

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major defects in the deposited Y2O3 films[33].

In this work, as –80 V bias and 4 sccm O2 flow rate are used to maintain

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low-energy ions bombardment and keep a metal sputtering mode, respectively, Y2O3

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thin films with different oxygen vacancy concentration have been achieved through DC reactive magnetron sputtering by varying Ts, and the oxygen vacancies dependent phase transition and hardness have been explored by X-ray photoelectron spectroscopy (XPS), X-ray diffraction (XRD), high-resolution transmission electron microscopy (HRTEM), and nanoindenter. 2. Experimental

2.1 Sample preparation Y2O3 thin films with a thickness of about 1 µm were grown on Si (001) wafer substrates by DC reactive magnetron sputtering a metal yttrium target (99.5 purity, 60 mm in diameter and 3 mm in thickness) in the mixed discharge gases of Ar (99.99%)

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and O2(99.99%) for 3 hours. Si (001) substrates were cleaned consecutively with acetone, alcohol and distilled water in an ultrasonic bath, and then were dried with nitrogen before introducing them into the vacuum chamber. After this, substrates were

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place at holder right above the target, keeping a 60 mm distance from yttrium target. Prior to deposition, the chamber was evacuated to a base pressure of 4×10-4 Pa by

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turbomolecular pump. During the deposition, the flow rates of Ar and O2 were

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controlled at 70 and 4 sccm, respectively, maintaining a metal sputtering mode to obtain relatively high deposition rate and stable sputtering. Various Ts ranging from

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room temperature (RT) to 600 ˚C were used to deposit Y2O3 thin films with different

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oxygen vacancy concentration, while substrate bias, deposition pressure, and sputtering current were kept at –80 V, 0.8 Pa, and 0.3 A, respectively. In addition,

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vacuum annealing treatments were performed on all samples at 1000 ˚C for one hour

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to study their structural stability, in which vacuum was below 10-4 Pa. 2.2 Characterization

The phase structures were determined by XRD using a Bragg-Brentano diffractometer (D8_tools) in θ-2θ configuration with Cu Kα radiation. Further observations on microstructure of the films were also performed using a field emission HRTEM (JEOL 2100F) operated at 200 kV. The core-level spectra of the films were detected by XPS (ESCALAB-250) spectrometer with a hemisphere detector (an energy resolution of 0.1eV supplied by an Al Kα radiation source), in which Ar+ cleaning procedure lasting 300s was applied to all samples to remove possible adventitious carbon and absorbed oxygen from the sample surface. The

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atomic concentration ratios of elements on the sample surface were calculated from the integral photoelectron peak intensities and known atomic sensitivity factors (ASFs)[39]. The hardness and modulus of films were measured using Nanoindenter

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(MTS XP) with continuous stiffness measurements (CSM) mode. A Berkovitch-type pyramidal diamond tip indented the films to a maximum depth of 800 nm. Constant

cr

stiffness data measurements were obtained by oscillating the tip during indentation

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with a frequency of 45 Hz and amplitude of few nanometers, and hardness values were taken at approximately 50-100 nm depth to avoid influence of the surface

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roughness and the substrate. At least six indentations at different places on the film

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surface were made. The curvatures of the substrates before and after deposition were measured by a surface profiler (Veeco Dektak 150), and the residual stress was

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calculated using Stoney equation[40]. In this work, the intrinsic stress was further

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obtained through the total measured residual stress by subtracting the thermal stress. 3． ．Results and discussion

3.1. Chemical composition and bonding XPS has been employed to calculate chemical composition of Y2O3 films by dividing integrated peak intensity for each element by corresponding sensitivity factor. The raw data was fitted using a linear combination of Lorentzian and Gaussian line shape (pseudo-Voigt-profile). It was found that Y-OH in Y spectra and OH-Y in O spectra were observed at the surface of Y2O3 due to hydroxylation caused by adsorption reaction of water, which may induce the variation of surface composition. In order to accurately calculate the O/Y ratio in lattice for Y2O3, we exclude the

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contribution of Y-OH in Y spectra and OH-Y in O spectra. Accordingly, O/Y atomic ratio was defined by the equation as follow. O/Y ratio =

IO-Y IY-O+IY 2 O 3 -x / SO1s SY 3 d

(1)

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The ‘I’ represents the integral intensity of corresponding peak and ‘S’ is the sensitivity factor of yttrium or oxygen. The calculated O/Y ratios as a function of Ts

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are shown in fig. 1 (a). Auger electron spectroscopy (AES) depth profile has also

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performed on film deposited at Ts = 600 ˚C to confirm the accuracy of the calculated method, which is shown in fig. 1 (b). The O/Y ratio obtained in the stable region is

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1.33, according with the XPS results.

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O/Y ratio for film grown at Ts = 25 ˚C (room temperature) is 1.52, which is slightly higher than the stoichiometry (1.5), confirming the complete oxidation of Y2O3 films.

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As Ts is increased to 200 ˚C, O/Y ratio drops to 1.45, turning into substoichiometry,

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and the oxygen concentration continuously decreases with further increasing Ts, implying a continuous increase in oxygen vacancies since oxygen vacancies are considered as the primary defects in substoichiometric Y2O3 films[41]. The decrease of oxygen concentration and deviation from stoichiometry in thin Y2O3 films by increasing Ts have also been reported by other researchers[24], which has been ascribed to that oxygen diffuses to interfacial area and reacts with substrate to form interfacial layer at high Ts. But this interfacial reactive effect should be ruled out due to much larger thickness (~1µm) in this work. During reactive sputter deposition of Y2O3 films, the relative O incorporation in the as-deposited films mainly results from the chemisorption of oxygen ions and atomic O generated in the plasma, dissociative

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chemisorption of O2, direct implantation, and recoil implantation. The raise of Ts can enhance dissociative chemisorption of O2, increasing oxygen concentration in Y2O3 films, on the other hand, Ts would increase the desorption probability of absorbed

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oxygen on growing surface, decreasing oxygen concentration. The later effect comes to play a dominant role in our experiment, resulting in a continuous drop of oxygen

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investigation on oxygen vacancies dependent phase transition.

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concentration with Ts and achievement of different oxygen vacancies suitable for

Fig. 2 shows the XPS spectra of Y 3d and O 1s of as-deposited Y2O3 films

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deposited at different Ts. Broadness and non-symmetry of the Y 3d spectra clearly

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indicate that there are more than one yttrium species presented in the films. Y 3d peak for Y2O3 film deposited at Ts = 25 ˚C can be deconvoluted into two different doublet

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peaks corresponding to the Y 3d5/2 and Y 3d3/2 electrons, revealing two different

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Yttrium bonding. The first doublet peaks located at 156.7±0.2eV and 158.7±0.2eV can be ascribed to Y–O bond[42, 43]; the second ones located at 158.4±0.2eV and 160.4±0.2eV are from hydroxylated Y-OH bond caused by atmospheric moisture[29, 44]. Although surface cleaning of the sample has been performed by Ar+ etching for 300s, the peak intensity from Y-OH bond is still strong, indicating that hydroxylated Y2O3 surface can easily form because of its lower heat of formation (−1435 kJ/mol per Y) as compared with the oxide (−953 kJ/mol per Y)[24]. In the O 1s spectra, two strong peaks located at 528.9±0.2 eV and 531±0.2eV appear, corresponding to the O–Y[26, 44] and O-H[44] bonds, respectively, further confirming the results of Y 3d spectra. In addition, a slight peak at 531.8±0.2eV is also found, attributing to O-C

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bond formed by little adventitious carbon[27, 45]. The rise of oxygen vacancies gives rise to that the O-Y peak for O 1s spectra shift slightly toward higher value by increasing Ts from 25 ˚C to 600 ˚C. However, Y 3d spectra exhibits a two-step

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variation with increasing Ts. As Ts is increased from 25 ˚C to 400 ˚C, Y-OH binding energies almost remain unchanged and the doublet peaks corresponding to Y–O bond

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continuously shift toward lower binding energies. Increment of oxygen vacancies can

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be responsibility for the observed low-energy shift of Y–O bond, which is in accordance with other reports[46]. Further increasing Ts to 600 ˚C, besides these

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peaks corresponding to Y–O and Y-OH bonds, a new doublet peaks appears at lower

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values, indicating that some amount of reduced oxide (Y2O3−x) occurs in the film[27, 47, 48]. The lower doublet peaks for Y 3d spectra have also been observed in reduced

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YSZ samples[49]. Therefore, critical oxygen vacancy contents exist in Y2O3 films,

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below which increment of oxygen vacancies just induces the low-energy shift of Y–O bond, above which the new reduced (Y2O3−x) oxide appears at lower binding energy due to formation of new yttrium surrounding with more oxygen vacancies. 3.2. Crystal structure and thermal stability XRD and HRTEM have been further performed to investigate the oxygen vacancies dependent phase structures and their stability. Fig. 3 displays the XRD patterns for both as-deposited and post-annealed Y2O3 films as a function of Ts. Firstly we focus on Ts dependent phase structures for as-deposited Y2O3 films. The XRD curve of Y2O3 films grown at 25 ˚C exhibit a broad peak with very low intensity centered at about 29.1˚, which is assigned to (222) of C-type cubic phase (ICDD file no. 41-1105),

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indicating a poor crystallinity. At Ts=200 ˚C, only a strong peak corresponding to (111) of B-type monoclinic phase (ICDD file no. 44-0399) appears, showing a strong (111) texture and good crystallinity. Increasing Ts to 400 ˚C, the monoclinic (111) peak

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almost disappears, and new non-symmetric peaks consisted of cubic (400) and monoclinic (11 2 ) occur, implying the presence of mixed cubic and monoclinic

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phases. At 600 ˚C, mixed phases of cubic and monoclinic still coexist in Y2O3 films,

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but preferred orientations turn to cubic (222) as well as monoclinic (401) and(40 2). Nevertheless, many small peaks corresponding to monoclinic phase also appear which

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results from large amount of random oriented monoclinic nanograins. Furthermore,

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noting the full width at half-maximum (FWHM) of main peaks for all as-deposited films, the maximum value of grain size appears in Y2O3 films deposited at Ts=200 ˚C,

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and competitive growth between cubic and monoclinic grains can be responsibility for

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the drop of grain size at higher Ts (400 and 600 ˚C). The detailed structural characters for as-deposited Y2O3 films deposited at Ts=600 ˚C have been further demonstrated by HRTEM, which is shown in fig. 4 (a). The film exhibits a number of nanoscale cubic and monoclinic domains as marked by ‘‘C’’ and ‘‘M’’, identifying by the corresponding fast Fourier transformations (FFT). Although the crystallite size for both cubic and monoclinic phases is smaller than 10 nm, some planar defects are still observed, which has also been reported in epitaxial Y2O3 films when the aggregates of oxygen vacancies exceed the critical value [38,50]. Meanwhile, some disordered regions, i.e., a loss of long-range order, as marked by ‘‘D’’, are present in the vicinity of the cubic and monoclinic domains, which is the

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results of competitive growth between cubic and monoclinic grains. In brief, the evolution of phase structure with Ts mainly includes that cubic phase presented at 25 ˚C transforms to monoclinic phase at 200 ˚C, and then changes to a

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mixture of cubic and monoclinic phases at higher Ts (400 ˚C, 600 ˚C), which should be dominated by variation of oxygen vacancy contents. By increasing Ts, different

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result that the transformation from monoclinic to cubic structure have been reported

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by Cho et al.[34], wherein the monoclinic phase just occurs at Ts≤400 ˚C when O/Y ratio is below 1.458, and phase transition from monoclinic to cubic takes place at

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Ts=500 ˚C while O/Y ratio approaches the theoretical ratio of 1.5, indicating that

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oxygen vacancies are benefit for the growth of monoclinic grains. Gaboriaud et al. [38, 51, 52] observed that large amount of oxygen vacancies could be introduced into

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epitaxial Y2O3 films by means of ions bombardment both during and after film

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growths, inducing the nucleation of monoclinic phase from cubic structure. They further put forward the model of extended defect nucleation that accumulation of oxygen vacancies in the {111} planes leads to a collapse of the lattice along the 〈111〉 direction, followed by a crystallographic shear along the 〈211〉 axes, consequently forming a new stacking sequence corresponding to a nucleus of the monoclinic phase. The observed planar defects caused by accumulation of oxygen vacancies shown in fig. 4 (a) imply that high concentration of oxygen vacancies promote the nucleation and growth of monoclinic phase even at Ts=600 ˚C. In addition, the following vacuum annealing experiment shows that monoclinic phase with oxygen deficiency is not thermodynamic stable at high temperature. Accordingly, high Ts also leads to the

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instability of monoclinic nucleus with oxygen deficiency, which is evidenced by HRTEM result that a large number of cubic nanograins also appear in film deposited at Ts=600 ˚C. Therefore, at relatively low Ts (200 ˚C), oxygen vacancies can induce

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the nucleation and growth of monoclinic phase and low Ts is not enough to induce the instability of monoclinic phase, thus only monoclinic phase occurs. At higher Ts (400

cr

˚C, 600 ˚C), the instability of monoclinic phase induced by high Ts begins to take

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effect, and the interaction of above two factors results in the competitive growth between cubic and monoclinic grains, forming the random-oriented finegrains

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evidenced by fig. 4 (a).

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Post annealing treatments for Y2O3 films have been mainly performed in the oxygen or air ambient[11, 34, 52, 53], and extra oxygen would be introduced to

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occupy the oxygen vacancies, giving rising to the transformation from monoclinic to

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cubic structure. Understanding the oxygen vacancies dependent stability of phase structure is also important for application as interface diffusion and reaction barrier coatings, however, less attention has been paid on. Consequently, vacuum annealing treatments at 1000 ˚C were applied to all samples to investigate the oxygen vacancies dependent structural stability. The O/Y ratio for film deposited at Ts=600 ˚C after vacuum annealing treatments has also been characterized by XPS, which remains almost unchanged as compared with as-deposited one. It is same for all as-deposited films from XRD results that annealing treatment at 1000 ˚C leads to narrowing of FWHM and shifting of peaks toward higher angle closer to standard peak position, corresponding to grain growth and relaxation of compressive stress, respectively,

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which are confirmed by following HRTEM and stress results. In addition, vacuum annealing also induce the structural variation. For film deposited at Ts = 25 ˚C, cubic structure remains stable after post-annealing, and the

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observed new peaks result from the recrystallization and growth of random-oriented grains from former (222)-oriented nanograins. For film deposited at Ts=200 ˚C, a

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remarkably strong monoclinic (111) peak still exist and the monoclinic structure

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remain stable even at 1000 ˚C, meanwhile some small peaks attributing to cubic structure appear, which may come from the transition from partial small-size

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monoclinic grains. For film fabricated at Ts=400 ˚C, strong monoclinic(11 2 )and cubic

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(400) peaks becomes one cubic (400) peak, meaning that phase transition from monoclinic to cubic takes place during high-temperature annealing. Besides, a small

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new peak at about 31.0˚ is observed, corresponding to neither monoclinic nor cubic

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phase, which can be ascribed to (211) of the oxygen defective phase (ICDD file no. 39-1063) [26,54]. For film with higher oxygen vacancies grown at Ts=600 ˚C, almost all monoclinic peaks disappear, and two strong peaks corresponding to cubic and oxygen defective phases occur.

HRTEM test has been carried out on post-annealed Y2O3 film deposited at Ts=600 ˚C to observe the detailed structural evolution. After post-annealed treatment, the random oriented cubic and monoclinic nanograins as well as disordered region for as-deposited sample begin to recrystallize, grow and transform to the observed lattice fringes along same direction, as shown in fig. 4 (b). The local FFT has been performed on A, B and C areas, corresponding to cubic phase, mixed phase of cubic

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and oxygen defective, and oxygen defective phase, respectively, which are consistent with the XRD result that monoclinic phase transforms to oxygen defective phase during annealing treatment for film with high concentration of oxygen vacancies.

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Furthermore, it is noteworthy that the post-annealed sample forms a coherent interface between cubic and oxygen defective phases with crystallographic relationship of

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(222)//(00 3 ) identified by FFT, meanwhile, a large amount of misfit dislocation

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appears at the interface region of two phases due to lattice mismatch. In order to further confirm that the oxygen vacancies dominate the phase transition from

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monoclinic to oxygen defective phase, a two-step annealing experiments, firstly air

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annealing treatment at 600 ºC for 8 h and then vacuum annealing at 1000 ºC for 1 h, were also carried out on film deposited at Ts=600 ˚C, and the corresponding XRD

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patterns have been given in fig. 5. All peaks shifted towards high angle after air

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annealing treatment at 600 ºC for 8 h, showing that the oxygen vacancies were filled by O from air. After the following vacuum annealing at 1000 ºC for 1 h, only cubic phase was observed. That means that high oxygen vacancies are crucial to form oxygen defective phase during post-annealing treatment. The combined results of XPS, XRD and HRTEM show that critical oxygen vacancy contents exist in Y2O3 films determining which phase structure the monoclinic phase transforms to during post annealing treatment, below which monoclinic phase transition to cubic phase for film deposited at Ts=400 ˚C, and above which transformation to the oxygen defective phase for film grown at Ts=600 ˚C. In spite of oxygen vacancy contents keep unchanged, oxygen defective phase just occurs in film

Page 16 of 36

after high-temperature annealing treatment, thus a big energy barrier for the formation of oxygen defective phase by means of rearrangement of high concentration of oxygen vacancies is required to overcome. Furthermore, although monoclinic

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structure for film deposited at Ts=200 ˚C can still keep stable after annealing treatment at 1000 ˚C, further increasing annealing temperature to 1100 ˚C results in a

cr

complete transition from monoclinic to cubic structure. In conclusion, B-type

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monoclinic phase with oxygen deficiency present in as-deposited Y2O3 films is not thermodynamic stable at high temperature, which can transfer to more stable phases

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including cubic or oxygen defective phase dependent on oxygen vacancy content. The

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higher transition temperature for monoclinic phase was observed in film deposited at Ts=200 ˚C, which owns low oxygen vacancy content, bigger grain size and stronger

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temperature of phase transition.

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texture, thus these factors would act together on the energy barrier and transition

3.3. Intrinsic stress and hardness

Fig. 6 illustrates the intrinsic stress for as-deposited and post-annealed Y2O3 films as a function of Ts. All as-deposited films exhibit compressive stress state, that first increases and then decreases with increasing Ts, and maximum value occurs at Ts=400 ˚C. The intrinsic compressive stress in the films originates from the ‘‘atomic peening’’ because -80 V bias were applied to substrate and growth surface suffered from the bombardment of incoming ions or atoms. The bombarding ions or atoms can implant into the subsurface as entrapped atoms and introduce high density of defects by collision cascade, further dilating the lattice and inducing compressive stress, which is

Page 17 of 36

consistent with the XRD peak shift. The increment of compressive stress with increasing Ts from 25 ˚C to 400 ˚C is mainly associated with presence of higher density of defects caused by the rise of oxygen vacancy content. The higher Ts (600

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˚C) remarkably enhances the migration ability of surface atoms and annihilation ability of Frankel pairs, and reduces the atomic peening effect even at more oxygen

cr

vacancies, thereby giving rise to drop of compressive stress.

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As compared with as-deposited films, annealing treatment results in obvious relief of compressive stress for films deposited at Ts=25 ˚C, 200 ˚C and 400 ˚C, and even

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stress reversal from compressive to tensile for film deposited at Ts=600 ˚C. Grains

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growth and phase transition took place during the annealed process, as evidenced by XRD and HRTEM, which would play an important role on stress variation. For films

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grown at Ts=25 ˚C and 200 ˚C, phase transition can be ruled out, and grains growth

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dominated the relief of compressive stress through film horizontal contraction caused by reduction of grain boundaries[11]. For films deposited at Ts=400 ˚C and 600 ˚C, phase transition should be also considered besides grains growth. Phase transition from monoclinic to cubic appears in film deposited at Ts=400 ˚C, which is associated with a slight decrease of ~1% volume since the theoretical densities of the cubic and monoclinic Y2O3 polymorphs are 5.030 g cm-3 and 4.980 g cm-3, respectively[14, 15]. The volume decrease generated by monoclinic→cubic phase transition can partially contribute to the stress relief. When phase transition from monoclinic to oxygen defective phase occurs in film deposited at 600 ˚C, volume would shrink due to that oxygen defective phase owns much higher density than B-type monoclinic structure.

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Phase transition of monoclinic→oxygen defective and grains growth act together on horizontal contraction, resulting in stress reversal from compressive to tensile. Fig. 7 displays the dependence of the hardness and modulus for as-deposited and

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post-annealed Y2O3 films on Ts. For as-deposited films, the hardness monotonously increases from 4.5 to 15 GPa with increasing Ts from 25 ˚C to 600 ˚C. The modulus

cr

exhibits the same trend with hardness. The hardness of bulk Y2O3 ceramic was

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reported in the range of 6.9 to 7.6 Gpa[55], which is consistent with the calculated hardness (7.7 GPa) by means of the dielectric chemical bond theory[56]. However,

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the reported hardness in Y2O3 films varied in a larger range of 6 to 12 Gpa[57],

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because more factors took effect in film, including defects, microstructure, texture, residual stress, grain size and phase configuration. In this work, the change of

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hardness has a same trend with variation of oxygen vacancy content. However, it has

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been reported that the increase of oxygen vacancy or reduction of the Y–O bonds can slightly reduce the hardness of Y2O3 through increasing the bond length and decreasing the bond strength[37, 58]. Thus the improvement of hardness with Ts could not be ascribed to increment of oxygen vacancy. The effects of other factors should be explored.

Lei et al.[57] have reported that columnar structures can be formed in Y2O3 films deposited under metallic-mode condition, leading to higher hardness, compared with non-columnar structures. Therefore, cross-section SEM investigations have been performed on as-deposited films, and the images are shown in Fig.8, in which typical columnar structure only appears at Ts = 200 °C. The observed poor crystallinity (Ts =

Page 19 of 36

25 °C) and the competitive growth between cubic and monoclinic grains (Ts = 400 °C and 600 °C) could suppress the columnar growth. Accordingly, in this work columnar structure does not play a dominant role in hardness enhancement.

ip t

In addition, in Y2O3 films, the cubic phase shows higher hardness than monoclinic structure[31] and (222) texture owns superior mechanical properties[57]. The lowest

cr

hardness at Ts = 25 ˚C is due to very poor crystallinity. At Ts=200 ˚C, the monoclinic

us

phase with good crystallinity shows higher hardness. The further improvement in hardness for films deposited at Ts=400 and 600 ˚C is caused by increase of cubic

an

phase and the presence of (222) texture. Moreover, stress hardening and grain size

and 600 ˚C.

M

strengthening should also partially contribute to the increment of hardness at Ts=400

d

After annealing treatment, the grains growth and stress relief take place, compared

Ac ce pt e

with corresponding as-deposited Y2O3 films, resulting in the reduce of stress hardening and grain size strengthening for post-annealed films. In fact, except for Y2O3 film fabricated at Ts = 25 °C, annealing treatment exactly leads to drop in hardness for Y2O3 films fabricated at Ts=200, 400 and 600 ˚C. The as-deposited Y2O3 film grown at Ts = 25 °C exhibits very low hardness due to very poor crystallinity or amorphous structure. Accordingly, for Y2O3 film fabricated at Ts = 25 °C, the improvement of crystallinity under annealing treatment gives rise to hardness enhancement, compared with as-deposited film. However the trend of the variation of residual stress and grain size with Ts is not always in accord with that of the variation of hardness with Ts for post-annealed films. This means that other factors also affect

Page 20 of 36

the evolution of hardness with Ts. Furthermore, it is noteworthy that high hardness just appears in Y2O3 films with mixed phase configurations including cubic + monoclinic for as-deposited films

ip t

deposited at Ts= 400 and 600 ˚C as well as cubic + oxygen defective phase for post-annealed film deposited at Ts= 600 ˚C. As cubic + monoclinic phases for

cr

as-deposited film deposited at Ts= 400 ˚C transform to cubic phase during annealing

us

treatment, the hardness drops remarkably and reaches the minimum value. For film with mixed phase configuration, the interface between different phases may block the

an

dislocation movement and give rise to the enhanced hardness[59]. Consequently,

M

Y2O3 film with high concentration of oxygen vacancies can maintain high hard even at the high temperature of 1000 ˚C by forming a mixed phases of cubic + oxygen

Ac ce pt e

4. Conclusions

d

defective, which would be benefit for some high-temperature applications.

Y2O3 thin films with different oxygen vacancy contents have been deposited by

reactive sputtering through varying Ts. O/Y atomic ratio decreases continuously from 1.52 to 1.37 with increasing Ts from 25 ˚C to 600°C, corresponding to an increase in oxygen vacancy concentrations. For as-deposited Y2O3 films, oxygen vacancies present at high Ts can promote the nucleation of monoclinic phase, whereas high Ts can also induce the instability of monoclinic phase. Consequently, the observed structural evolution is dominated by the competition between formation of monoclinic phase caused by oxygen vacancy concentrations and instability of monoclinic phase induced by high Ts, in which cubic phase present at 25 ˚C transforms to monoclinic

Page 21 of 36

phase at 200 ˚C, and then changes to a mixture of cubic and monoclinic phases at higher Ts (400 ˚C, 600 ˚C). During post annealing treatment, the cubic structure goes on keeping stable, whereas the monoclinic phase with oxygen deficiency present in

ip t

as-deposited Y2O3 films is not thermodynamic stable at high temperature, which can transfer to more stable phases including cubic or oxygen defective phase depending

cr

on oxygen vacancy contents. There is a critical value of oxygen vacancy content,

us

below which monoclinic phase transits to cubic phase for film deposited at Ts=400 ˚C, and above which transforms to the oxygen defective phase for film grown at Ts=600

an

˚C. Annealing treatment also leads to grain growth and relaxation of compressive

M

stress, further reducing the hardness by weakening stress hardening and grain size strengthening. Compressive stress releases mainly through film horizontal contraction

d

caused by grains growth and phase transition. The high hardness just appears in Y2O3

Ac ce pt e

films with mixed phase configuration including cubic + monoclinic for as-deposited films deposited at Ts= 400 and 600 ˚C as well as cubic + oxygen defective phase for post-annealed film deposited at Ts= 600 ˚C, in which the interfaces between different phases may block the dislocation movement and give rise to the enhanced hardness. To maintain stable as well as favorable mechanical performance in a large temperature range, oxygen vacancies could be introduced to help maintaining multiphase structure in Y2O3 films.

Acknowledgements The support from National Natural Science Foundation of China (Grant Nos.

Page 22 of 36

51672101, 51602122, 51102111, 51572104 and 51372095), the National Key Research and Development Program of China (2016YFA0200400), the Aviation Science

Foundation

of

China

(No.201430R4001), the

NSF

of

Jilin

ip t

Province (No.20160520010JH) China, National Major Project for Research on

cr

Scientific Instruments of China (2012YQ240264), is highly appreciated.

us

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*Figure Captions Fig. 1. (a) The calculated O/Y ratios of as-deposited samples from XPS. (b) The typical Auger electron spectroscopy depth profile for Y2O3 film deposited at Ts = 600

ip t

˚C.

Fig. 2. XPS core level spectra of (a) Y 3d and (a) O 1s for Y2O3 films deposited at

cr

different Ts.

us

Fig. 3. (a) XRD patterns both for as-deposited and post-annealed Y2O3 films as a function of Ts. (b) Enlarged XRD peaks to detailed observation of main characteristic

an

peaks ranging from 25˚ to 35˚.

M

Fig. 4. (a) HRTEM image of the as-deposited sample grown at Ts=600 ˚C: The cubic, monoclinic and disordered regions are designated by ‘‘C’’, ‘‘M’’ and ‘‘D’’,

d

respectively, and the inserted FFT images correspond to the domains indicated by the

Ac ce pt e

white lines. (b) HRTEM image of the post-annealed sample grown at Ts=600 ˚C: The insets correspond to the local FFT performed on different areas indicated by white box.

Fig. 5. XRD patterns of as-deposited and different post-annealed Y2O3 film fabricated at Ts=600 ˚C. Oxygen defective phase no longer forms during a two-step annealing treatment (marked by a→c→d, firstly exposed in air at 600 ºC for 8 h and then annealed in vacuum at 1000 ºC for 1 h), while oxygen defective phase appears through just vacuum annealing treatment at 1000 ºC for 1 h (marked by a→b).

Fig. 6. Intrinsic stress for as-deposited and post-annealed Y2O3 films as a function of Ts.

Page 27 of 36

Fig. 7. Hardness and modulus for as-deposited and post-annealed Y2O3 films as a function of Ts.

Fig. 8. Cross section SEM images of the as-deposited films deposited at Ts = 25 °C

Ac ce pt e

d

M

an

us

cr

ip t

(a), 200 °C (b), 400 °C (c) and 600 °C (d).

Page 28 of 36

ip t cr

Ac ce pt e

d

M

an

us

Fig. 1.

Page 29 of 36

ip t cr us an M d Ac ce pt e

Fig. 2.

Page 30 of 36

ip t cr us an

Ac ce pt e

d

M

Fig. 3.

Page 31 of 36

ip t cr us an M d Ac ce pt e

Fig. 4.

Page 32 of 36

ip t

Ac ce pt e

d

M

an

us

cr

Fig. 5.

Page 33 of 36

ip t cr us an

Ac ce pt e

d

M

Fig. 6.

Page 34 of 36

ip t cr us an M d Ac ce pt e Fig. 7.

Page 35 of 36

ip t cr us an

Ac ce pt e

d

M

Fig. 8.

Page 36 of 36