Tungsten Incorporation into Gallium Oxide - American Chemical Society

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Tungsten Incorporation into Gallium Oxide: Crystal Structure, Surface and Interface Chemistry, Thermal Stability and Interdiffusion Ernesto J. Rubio, Thomas E. Mates, Sandeep Manandhar, Manjula I. Nandasiri, Vaithiyalingam Shutthanandan, and Chintalapalle V Ramana J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b05487 • Publication Date (Web): 14 Oct 2016 Downloaded from http://pubs.acs.org on October 18, 2016

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The Journal of Physical Chemistry

Tungsten Incorporation into Gallium Oxide: Crystal Structure, Surface and Interface Chemistry, Thermal Stability and Interdiffusion E.J. Rubio1, T. E. Mates2, S. Manandhar1,3, M. Nandasiri3, V. Shutthanandan3 and C.V. Ramana1∗

1

Department of Mechanical Engineering,

University of Texas at El Paso, El Paso, Texas 79968, USA 2

Materials Department, University of California, Santa Barbara, California 93106, USA 3

Environmental Molecular Sciences Laboratory (EMSL), Pacific Northwest National Laboratory (PNNL), Richland, Washington 99352, USA

∗

Author to whom correspondence should be addressed; Email: [email protected];

Tel: 915-747-8690; Fax: 915-747-5015 1 ACS Paragon Plus Environment


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ABSTRACT Tungsten (W) incorporated gallium oxide (Ga2O3) (GWO) thin films were deposited by radio-frequency magnetron co-sputtering of W-metal and Ga2O3-ceramic targets. Films were produced by varying sputtering power applied to the W-target in order to achieve variable W-content (0-12 at%) into Ga2O3 while substrate temperature was kept constant at 500 °C. Chemical composition, chemical valence states, microstructure and crystal structure of as-deposited and annealed GWO films were evaluated as a function of W-content. The structural and chemical analyses indicate that the samples deposited without any W-incorporation are stoichiometric, nanocrystalline Ga2O3 films, which crystallize in β-phase monoclinic structure. While GWO films also crystallize in monoclinic β-Ga2O3 phase, W-incorporation induces surface amorphization as revealed by structural studies. The chemical valence state of Ga ions probed by X-ray photoelectron spectroscopic (XPS) analyses is characterized by the highest oxidation state i.e., Ga3+. No changes in Ga chemical state are noted for variable W-incorporation in the range of 0-12 at%. Rutherford backscattering spectrometry (RBS) analyses indicate the uniform distribution of W-content in the GWO films. However, XPS analyses indicate the formation of mixed valence states for W ions, which may be responsible for surface amorphization in GWO films. GWO films were stable up to 900 oC, at which point thermally induced secondary phase (W-oxide) formation was observed. A transition to mesoporous structure coupled with W interdiffusion occurs due to thermal annealing as derived from the chemical analyses at the GWO films’ surface as well as depthprofiling towards the GWO-Si interface. Surface imaging analyses indicate thermally induced morphological changes are dependent on W-concentration in the GWO films.

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Thermally induced diffusion of W in the film is responsible for the observed formation of pores of variable size; the maximum pore radius noted was ~27 nm for GWO films with highest W-content. The electronic charge redistribution appears to be dominated by the hydroxyl groups and W-chemistry as evident in XPS analyses. RBS data indicate that the extent of diffusion and intermixing layer depth is dependent on W-content in the GWO films. Thermally induced W-diffusion and depth penetration into the Si substrate with SiW-Ga2O3 intermixing at the interface is evident only in GWO samples with highest (12 at%) W-incorporation. A model has been formulated to account for the mechanism of Wincorporation, thermal stability and interdiffusion via pore formation in GWO films. Keywords:

Ga2O3 thin films; W-doping; Chemical Composition; Interdiffusion; RBS



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INTRODUCTION Interest in the wide band gap oxide materials is continuously increasing in recent years due to their wide range of technological applications in electronics, photonics, electro-optics, opto-electronics, photo-catalysis, chemical sensing, and magnetoelectronics. These materials find numerous applications in photo-resistors, photodiodes, luminescent phosphors, electronic switches, and gas sensors. Gallium oxide (Ga2O3), which is a stable oxide of Ga, is a wide band gap (~5 eV) material.1–4 Ga2O3 based thin films and nanostructures are attractive for application in electronics, luminescent phosphors, high-temperature sensors, antireflection coatings, lithium batteries and solar cells.1–19 Ga2O3 has been recognized as a deep ultraviolet transparent conducting oxide (UV−TCO), which makes the material a potential candidate for transparent electrode applications in UV optoelectronics, photonics and thin-film transistors.1–4,9–11,17 In addition, it has been reported and demonstrated in the literature that Ga2O3 presents interesting physico-chemical properties and device applications, which can be tuned by the presence of dopants. Catalytic activity for next generation reactors for hydrogen production, reduction of carbon dioxide, and decomposition of organic compounds are some of the most important ones to mention.20–25 However, in all these aforementioned applications, fundamental understanding of the physical chemistry of Ga-oxide based materials so as to provide a better control on the interplay between structure, thermodynamic conditions, chemical processes and kinetics is the key to achieve enhanced efficiency as well as control over the environmental pollutants. Currently, in addition to the single phase Ga-oxide nanostructured materials, Ga2O3-based multi-component architectures or hybrid materials with nanostructured


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morphologies are increasingly gaining attention in view of their novel functionalities or new applications.2,4,12,16–19 For instance, excellent luminescence was observed in Sn and Cr doped Ga2O3 nanowires.4,9 Patil et al. have recently reported on the utilization of Ga2O3-graphene oxide hybrid nanomaterials with enhanced electrochemical performance in lithium ion batteries.19 In their work, an approach is described to control the crystal phase of nanostructures by the hybridization of Ga2O3 with reduced graphene oxide (rGO).19 It was demonstrated that a nanomaterial with a mixed α-Ga2O3/β-Ga2O3/γ-Ga2O3 phase delivers the best performance, highlighting the importance of control over the chemistry and phase in optimizing the electrochemical activity.19 Similarly, solutionprocessed Ga2O3 dielectrics are promising to contribute the development of next generation, large area green-electronics.2 Solution-processed In2O3/Ga2O3 hybrid structures in thin-film transistors with the desired mobilities and negligible hysteresis are expected to open the new possibilities to the low-cost and large-area green oxide electronics.2 Enhanced photocatalytic activity for application in energy-harvesting devices has been reported for metal-incorporated Ga2O3 or metal/Ga2O3 hybrid materials.14,17,24,25 Doping metal ions, such as Ni, Zn and Pb, into Ga2O3 has been reported to enhance the photocatalytic activity of Ga2O3.17,24,25 Most recently, Zhang et al. have reported another approach based on liquid metal/Ga2O3 frameworks to serve as an efficient photocatalytic system.14 However, while evolving new functionalities and emerging applications involving Ga2O3 nanomaterials or hybrid structures is promising, it is very clear that the surface/interface chemistry and control over the phase and composition is the key to tailor the materials’ performance to meet the requirements of a given application or to search for new technological applications. Most importantly,



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developing a fundamental scientific understanding of the mechanisms and insights into how these complex structures are formed along with evaluating their thermodynamic stability, especially for metal ion incorporated Ga2O3, is very important. Gallium oxide exhibits five polymorphous: α, β, γ, δ, and ε. Among these polymorphs, β-Ga2O3 is the most stable phase, both chemically and thermodynamically, with a melting point of 1780 °C.26–30 The β-phase Ga2O3 adopts the monoclinic crystal structure with two distinct gallium sites (six- and four-coordinated) and three oxygen sites (three- and four-coordinated). 26–30 The lattice parameters of β-Ga2O3 are a=12.2140 Å, b=3.0371 Å, c=5.7981 Å, and β=103.83°.26–32 Intrinsic Ga2O3 is a wide band gap insulating oxide; but, it exhibits an n-type semiconductor behavior when selectively doped with certain type of metal ions or due to the presence of oxygen ion vacancies.30,31 However, while the mechanism of n-type conduction in β-Ga2O3 appears to be mostly by the generation of electrons from oxygen vacancies ionization, the overall electrical conductivity of the resulting materials upon doping is still under debate.30,31 Doping with carefully chosen metallic ions can significantly alter the properties of β-Ga2O3 making it suitable for a wide variety of technological applications. From optical properties point of view, doping with Mn, Si, Cr, and Cu metals into β-Ga2O3 has been reported to induce changes in the optical absorption and band gap, although the extent of changes depend on the nature of the metal.33–35 Recently, we have reported that tungsten (W) incorporation into Ga2O3 can induce a significant shift in the band gap.36 As reported previously, incorporation of ~11% at. of W during co-sputtering deposition was capable of introducing a shift in the absorption edge of gallium oxide by narrowing the band gap value from ~5.0 eV to ~4.3 eV ).36 A wide range of possible applications, in addition to


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improving the response time for high-temperature sensing, were predicted for GWO films with tunable absorption band edge and optical properties as a function of tungsten content. 36,37

However, a more detailed account of the effect of W-doping into Ga2O3 and a

correlation between W-concentration, associated chemistry and structural properties is totally missing at this time. On the other hand, a detailed account of surface/interface chemistry, electronic structure and thermo-chemical stability is quite important in order to understand the fundamental science so as to tune the properties of GWO films for enhanced performance in practical device applications. Furthermore, understanding the fundamental science of this W-Ga2O3 model system may shed light on the underlying mechanism, which could be useful and applicable to a large class of metal-incorporated Ga-oxide based materials. In this context, the present work was performed in order to derive a comprehensive understanding of the chemical composition, surface morphology, and structural properties for GWO films made by co-sputter deposition and with variable W-content. In addition, an attempt is also made to evaluate effect of post-deposition annealing on the chemistry and structure of GWO films to understand their thermochemical and structural stability, which is another important consideration for their feasibility into practical device applications, especially high-temperature sensors. For our work, while the ultimate goal is to realize novel materials with improved response time and sensitivity towards oxygen at elevated temperatures in chemical combustion systems, the impetus is towards the fundamental understanding of the effect of W-doping into Gaoxide nano-materials and thermal stability. As such, utmost attention is paid towards the surface and interface chemistry evaluation and thermally induced interfacial structure modifications in GWO nanocrystalline films using a set of well-targeted analytical



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methods. Specifically, combined X-ray photoelectron spectroscopy (XPS) and Rutherford backscattering spectrometry (RBS) together with microscopic measurements were employed to accurately probe the W-induced effects, thermally-induced surface and interface composition changes and to model the associated physical/chemical mechanism upon annealing. Based on the results as presented and discussed in this paper, a chemistry-structure-stability relationship is derived to account for the changes induced by W-content in addition to thermally-induced interfacial chemical changes in GWO nanocrystalline films. EXPERIMENTAL DETAILS A.

Samples’ Fabrication

Film Deposition. Tungsten incorporated Ga2O3 films were deposited onto silicon (Si) (100) wafers by radio-frequency sputtering. All the substrates were thoroughly cleaned and dried with nitrogen before introducing them into the vacuum chamber, which was then evacuated to a base pressure of ∼10-6 Torr. Deposition was made by co-sputtering of W-metal and Ga2O3 ceramic targets. Ga2O3 and W targets (Plasmaterials, Inc.) of 2 in. diameter and 99.999% purity were employed for every set of depositions made. Both targets were placed on a two different, 2 in. sputter guns, which were placed at a distance of 7 cm from the substrate. A sputtering power of 25 W was initially applied to each target while introducing high-purity argon (Ar) into the chamber to ignite the plasma. Once the plasma was ignited the power for each target was increased to their respective sputtering power for reactive deposition. The flow of Ar and oxygen (O2) were controlled using an MKS mass flow meters. Before each deposition, the targets were pre-sputtered for 20 min with a closed shutter above the gun. The deposition was carried out for 30 min.


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The thickness of the films was in the range of 55−32 (±3) nm. Deposition of GWO samples was carried by keeping Ga2O3 sputtering power constant at 100 W and varying the W-power in the range of 0-100 W. The intended purpose of variable sputtering power to W-target is to vary the amount of W-content into Ga2O3 films or produce GWO samples with variable W-content. The details of sputtering power and samples produced are listed in Table I. The substrate temperature (Ts) was fixed at 500 °C for all the films. The primary reason for selecting this deposition temperature is due to the fact that 500 °C is critical to promote the formation of nanocrystalline, β-Ga2O3 films, as reported previously elsewhere.26 For comparison, intrinsic Ga2O3 films without any W-dopant were also deposited under the same temperature and sputtering power conditions. Thermal Treatment. The GWO samples deposited with variable W-content were subjected to thermal treatment by annealing them at a high temperature (900 oC) to understand their thermal and chemical stability. After the evaluation in their as-deposited state, the GWO samples heat treated at 900 oC for 1 hour. The GWO films were introduced into a programmable furnace, where the constant temperature ramping rate was kept at 10 °C/minute to reach the final temperature of 900 oC. The samples were then set for one hour and then they were cooled down by natural convection. The structural, chemical and properties of the annealed GWO films were studied. B.

Characterization

XRD. Grazing incidence X-ray diffraction (GIXRD) measurements on GWO films deposited on Si were performed using a Bruker D8 Advance X-ray diffractometer. All the measurements were made ex situ as a function of variable W-content. XRD patterns were recorded using Cu Kα radiation (λ = 1.54056 Å) at room temperature.



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XPS. X-ray photoelectron spectroscopy (XPS) was used to corroborate composition of the films, and determine the chemical bonding of the ionic species in the film. Furthermore, XPS measurements allow probing the chemical characteristics of the surface, and any changes in composition compared to the bulk section of the films. XPS measurements were, therefore, made on the as-deposited GWO samples with variable Wcontent and also on the GWO films annealed at 900 °C for 1 hour. XPS surface analysis and depth profiling were carried out on a Kratos Axis Ultra DLD Spectrometer (Kratos Analytical, Manchester UK) which consists of a high performance Al Kα monochromatic x-ray source (1486.6 eV) and a high resolution spherical mirror analyzer. Survey scans were typically carried out at pass energies of 80 or 160 eV, while high-resolution scans were performed at a pass energy of 20 eV. For depth profiling the samples were sputtered using 5 KeV Ar+ ion beam rastered over a 2 × 2 mm2 area of the sample. The high resolution XPS spectra were recorded at pass energy of 20 eV with step size of 0.1 eV. The pass energy 20 eV in the 700 × 300 µm analysis area is referred to the full-width at half-maximum (FWHM) of 0.59 eV for Ag 3d5/2. The charge neutralizer with low energy electrons was used to exclude the surface charging effects and the binding energy of C 1s at 284.8 eV was used as the charge reference. XPS depth profiling data was analyzed by CasaXPS software using Gaussian/Lorentzian (GL(30)) line shape and Shirley background correction. For concentration of various elements present i.e., Ga, W, Si and O, the error of estimation is ±0.01 at% for metals (Ga, W, and Si) and ±0.1 at% for oxygen. RBS. Rutherford backscattering spectrometry (RBS) measurements were made using a NEC 3MV tandem accelerator (9SDH-2) to understand the chemical composition and


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elemental depth distribution of the GWO−Si (100) samples. An incident ion beam of 2 MeV He+ at 7° angle of incidence was used. The backscattered ions were detected using a silicon barrier detector at a scattering angle of 150°. Composition and depth profiles were determined by simulating the experimental backscattered spectra from ion beam analysis using SIMNRA software.38,39 The error of estimation is ±0.02 at% ±1 nm for concentration and thickness/depth values, respectively. AFM. Surface morphology of the GWO films was studied employing atomic force microscopy (AFM) using a Veeco Multimode scanning probe microscope with a Nanoscope V controller. A standard procedure40 for thin film surface characterization using AFM has been employed. AFM images were acquired using the ScanAsyst mode which utilizes a Bruker proprietary method for curve collection and sophisticated algorithms to continuously monitor image quality, and automatically make appropriate parameter adjustments. Aluminum coated silicon cantilevers (Bruker, USA) were used to acquire ScanAsyst mode images. The cantilevers measure 115 µm long, 25 µm wide and 0.65 µm thick with a spring constant of 0.4 N/m a resonance frequency of 70 kHz. The calibration standard used to calibrate the scanner consisted of platinum-coated, 200 nmtall silicon columns spaced at 10 µm intervals on centers. The columns have a length of 5 nm on a side. The drive amplitude varied between 30 and 180 mV. The images were then subjected to a 3rd order flattening procedure using the Veeco Nanoscope software to remove the non-linear background artefact introduced by the piezo scanner. SEM. Surface imaging analysis was performed using a high-performance and ultrahigh-resolution scanning electron microscope (Hitachi S-4800). Secondary electron



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imaging was performed on GWO films grown on Si wafers using carbon paste at the ends to avoid charging problems. RESULTS A. Surface and interface chemistry – RBS and XPS Chemical composition analysis of the GWO co-sputtered samples made using RBS measurements confirmed the presence of W in all the samples. The RBS spectra of GWO films are shown in Fig. 1. The data shown are for GWO films deposited under variable W-sputtering power. The experimental and simulation data are shown. The backscattered ions observed were due to various elements and their respective energy positions are as indicated by arrows for the experimental spectrum. The scattering from W, the heaviest among the elements present in either the film or substrate, occurs at higher backscattered energy (1780 keV). Similarly, for Ga, the peak is located at the backscatter energy of 1580 keV as shown in the RBS spectra (Fig. 1). The measured height and width of the respective peaks are related to the concentration and thickness distribution of Ga and W atoms in the GWO oxide film and serves as a calibration check for composition and thickness since known Rutherford scattering cross section and experimental parameters can be used to calculate the height and width.41,42 As indicated in Fig. 1, the step edge and peaks due to ion backscattering from Si (substrate) and O atoms (film) are observed at 1100 and 660 keV, respectively. The composition and thickness of the films were determined by simulating the experimental spectrum for the set of experimental conditions. The procedure utilized to derive the chemical composition of the grown films is as represented in Figure 1. The experimental curve (circles) along with the simulation curve (lines) calculated using SIMNRA code are shown in Fig. 1. The


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simulated curve was calculated using SIMNRA code38,39 for the fixed set of experimental parameters: (1) incident He+ ion energy, (2) integrated charge, (3) energy resolution of the detector, and (4) scattering geometry. It can be seen (Fig. 1) that the simulated curves (solid line) calculated using the experimental parameters are in good agreement with the experimental RBS spectra. This observation indicates that film composition reasonably simulates the spectra and provides the estimate of composition. The results obtained from the simulation of the RBS data indicate the increasing percentage of W in the films with increasing sputtering power. The data are presented in Table I, where film thickness variation with sputtering power is also presented. Increasing W-content from 0 to ~12 at% in the GWO films with increasing sputtering power from 0 to 100 W. RBS data coupled with spectroscopic ellipsometry analyses (not shown) thus provide a calibration and hence the means to control the amount of Wcontent and thickness of GWO films. Since the time of deposition is constant, film thickness increases with increasing W-content. X-ray photoelectron spectroscopy measurements also allowed us to determine the atomic percentage of tungsten as well as the chemical valence states of the constituent ions present in the GWO samples. The XPS survey spectra of representative GWO films are presented in Fig. 2. The XPS spectra indicate that Ga, W, and O are the main constituents of the deposited films. The presence of C 1s is evident in the survey spectra; the carbon peak in the XPS spectra is due to adventitious carbon from exposure to air following fabrication, before being placed in the XPS system. Therefore, the spectra were calibrated to the C 1s peak at a binding energy (BE) of 284.6 eV. However, sputtering with Ar+ ions fully eliminated the C 1s peak in the survey scans (not shown)



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which is a clear indication that the carbon is a result of adsorbed species on the film surface due to air exposure. The absence of a Si substrate peak is due to the fact that the deposited Ga-oxide films cover the Si surface and are sufficiently thick to prevent detection of photoelectrons originated from the substrate. In order to determine the chemical states of the metal ions present in the films, detailed analysis of the core level spectra of the main constituents was performed. The high resolution scans of Ga 3d, W 4f and O 1s were recorded and analyzed to obtain their chemical states and the interaction between respective metal ions and oxygen. The corelevel XPS spectra of Ga 3d region are shown in Fig. 3a. The data shown are for samples with a variable W-content. It can be seen that the Ga 3d peak is located at a binding energy (B.E.) of ~20.5 eV. The XPS Ga 3d core-level peak exhibits a doublet corresponding to Ga 3d5/2 and Ga 3d3/2 (Fig. 3a) at the binding energy values of 20.5 eV and 20.0 eV, respectively.43 The Ga 3d5/2 peak at 20.5 eV with a spin-orbit splitting energy (∆E=B.E(Ga 3d5/2 - Ga 3d3/2)) of 0.5 eV in this work is in good agreement with the literature values characterizing the Ga6+ state.9,44-46 In bulk Ga2O3, this Ga 3d peak appears at 20.6 eV.44,45 Thus, the BE location of Ga 3d peak at ~20.5 eV accounts for the Ga-O bonds, where Ga ions exist in their highest oxidation state i.e., Ga3+.44–46 Furthermore, the lower B.E. (~19 eV) components, which are characteristic of metallic Ga, were not seen in this work. For instance, for intrinsic and In-doped Ga2O3 nanostructures, the presence of a minor component at ~19 eV was attributed to a very small amount of either metallic Ga or formation of Ga sub-oxide on the surface.9 Absence of such features in this work indicate that, under the thermodynamic conditions employed for fabrication, all the GWO samples contain Ga in their fully oxidized state


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i.e., Ga6+. It is also interesting to note that the Ga 3d peak is constant without any appreciable change in either peak shape or BE position, regardless the amount of W incorporated in to the films. This is an indication that the W-ions incorporated into Ga2O3 do not induce any changes in the chemical state of gallium. The detailed core-level spectra of W 4f for GWO films are shown in Fig. 3b. The XPS W 4f core-level peak exhibits a well-resolved doublet corresponding to W 4f5/2 and W 4f7/2 at BE~37.8 eV and 35.7 eV, respectively. The W 4f7/2 peak at 35.7 eV is in agreement with the literature value of 35.7 eV characterizing the W6+ state in WO3.47–50 It is noted that the peak intensity increases with increasing sputtering power to the W-target and, thus, increasing W-content in the GWO films. Similar to Ga 3d peak, no shift in the BE of W 4f peaks is noted with increasing atomic percentage (up to 12 at%) of Wincorporated into Ga2O3. However, it is noted that W 4f doublet exhibits peak broadening, which is an indicative of the fact that some of the W ions exists in their lower valence states. The height of W 5d and W 4f peaks increase with increasing sputtering power for the W-target indicating an increase in the atomic percentage of tungsten in the GWO films. This result is in corroboration with the RBS data presented in the previous section. The high resolution core-level XPS spectra of O 1s are presented in Fig. 3c. The data shown are as a function of variable W-sputtering power. The peak position for low concentration of W is located at BE~530.8 eV, which corresponds to the position of O 1s for Ga2O3 when gallium is present in its highest chemical valence state (Ga+3).28,51,52 The O 1s is observed to shift towards slightly lower BE side with increasing W-atomic concentration. Higher content of W in the GWO films shift the peak to BE~530.6 eV, which is reported to be the characteristic of tungsten bonded with oxygen.49,50 Note that



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the O 1s peak is broader in the as-deposited GWO samples. In addition to the main contributions that can be attributed to oxygen bonded to Ga and W, there is a small shoulder appears at BE~532.3 eV. This can be attributed to oxygen bonded to carbon, and possibly in part to surface hydroxyl groups.15,28 The presence of surface hydroxyl groups can be attributed to the sample exposure to air following fabrication, before being placed in the XPS system. Most important to note is the fact that the overall concentration of W determined using RBS is in good agreement that probed using XPS. The chemical composition analyses, thus, indicate that the W-content vary in the GWO films from 0 to 12 at% for a variation of sputtering power applied to the W-target from 0 to 100 W. This atomic percentages correspond to the total W-content i.e., ionic and metallic percentage of W, inside the GWO film without considering the contributions from C or H2 adsorbed. The highest W-amount (12 at%) is recorded in GWO films when the sputtering power to the W-target was equal to that of Ga2O3 ceramic target i.e., 100 W (see, Table I). B. Surface structure – GIXRD The GIXRD patterns of the GWO films are presented in Fig. 4. The data are shown as a function of W-content in the films. It is evident that the peaks become diffuse with increasing W-content in GWO films which indicates W-induced amorphous nature in the GWO samples. Intrinsic Ga2O3 films are nanocrystalline as seen in the XRD patterns. Also, our previous results demonstrated the formation of crystalline GWO films under the variable W-content incorporation into Ga2O3 matrix; such films were crystallized in the corresponding β-Ga2O3 phase.36 However, for GWO films, the surface structure was probed using GIXRD measurements. For intrinsic Ga2O3 films, peaks appeared in the GIXRD pattern indicating the crystalline nature of the GWO surface


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layers and the peaks identified still corresponds to monoclinic β-phase. The (512) peak located at ~64.5° was found to be the most intense peak. The other peaks noted are (40 1 ) and (111) at ~40° and ~36.5°, respectively.52 The GIXRD patterns indicate very interesting results for W-incorporation into Ga-oxide matrix. While W-incorporation maintains the overall crystalline nature of the bulk of the samples, the surface layers become amorphous as is evident in GIXRD patterns, where the patterns become diffuse with increasing W-content. C. Surface morphology – SEM The SEM images indicating the surface morphology of intrinsic Ga2O3 and GWO films are presented in Fig. 5. SEM images demonstrate the characteristic feature of granular morphology for intrinsic Ga2O3 films with an average grain size of ~25 nm. This observation is in good agreement with previous studies, where crystallization of Ga2O3 films is noted at a deposition temperature of 500 °C.26 The change in the surface morphology of the GWO films as a result of tungsten incorporation is evident in Fig. 5. The film surfaces become featureless with increasing W-content. The surface of the Wdoped films become smooth without any presence of granular morphology that is seen for intrinsic Ga2O3 films, with segregated or occasional grains on the surface. This observation corroborates with the GIXRD results, where the formation of amorphous surface layers was evident for increasing W-content from 0 to 12 at% in GWO films. D. Thermal stability and interdiffusion The RBS data of annealed (900 oC) GWO films are shown in Fig. 6. The data shown are for a variable W-content. It can be seen in the RBS spectra that the samples with higher W-content experience some changes in the peak shape and height of Ga and 17 ACS Paragon Plus Environment


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W. However, there are no changes observed for intrinsic Ga-oxide and GWO samples deposited with a W-sputtering power in the range of 0-75 W. The most significant changes were noted for the samples deposited with 100 W sputtering power to the Wtarget i.e., the highest content (12 at%) of W in the Ga-oxide matrix. It can be seen in the RBS spectra of GWO samples that there is some degree of tailing or shouldering effect at the left of the peaks Ga and W peaks; W-peak changes are most dominant. Such tailing or shouldering of the peak towards higher energy indicates the diffusion occurring at the interfacial region of the GWO film and Si substrate. As seen in Fig. 6, this effect is much clear and dominant in W peak of GWO-100 samples. The interdiffusion of W is evident by tailing of the W peak at the interface coupled with some degree of Ga diffusion. In order to account for the observed changes and simulate the chemical structure changes, the RBS spectra were then fitted using SIMNRA by creating multiple layers with different composition of W, Ga, Si and O until the experimental data and SIMNRA calculated data match with each other, particularly the tailing at the interface. In other words, the depth profile fitting has been adopted to verify the chemical composition changes across the interface by fitting the data. Clearly, intrinsic Ga2O3 and GWO-50 samples had sharp interface between film and the substrate with no varied concentration of W at varying depth of the film. The GWO-75 film shows very small but almost negligible intermixing layer as shown by the depth profile compared to GWO-100 film, which exhibits significant intermixing layer. The depth of W diffusion depends upon the concentration of total W-content in the GWO film. For instance, for GWO-75, the depth profile indicates that the W atoms diffused 530.6 eV.49,50,54 Nevertheless, it is important to mention that O 1s peak remained at the same position for GWO-75 and GWO-100 due to the similar surface W-concentration. The presence of a peak broadening for O 1s indicates the existence of metal hydroxyl group, which is reasonable due to the exposure of the films to open atmosphere.55–57 The GIXRD spectra obtained for the GWO samples are shown in Fig. 10. It is evident that all the patterns exhibit peaks indicating the crystalline nature of the samples. In the case of intrinsic Ga2O3 films, data are similar before and after annealing indicating the structure stability in intrinsic β-Ga2O3 films. However, for GWO films, the evolution of diffraction peaks corresponding to both monoclinic Ga2O3 and monoclinic WO3 are evident indicating their crystalline transformation. The observed peaks can be indexed to monoclinic Ga2O3 (JCPDS No. 00-041-1103) and monoclinic WO3 (JCPDS No. 00-0431035). The very first implication of this observation is, perhaps, the formation Ga2O3WO3 mixed oxide upon thermal annealing. Also, similar to as-deposited GWO films, the peaks noted correspond to (40 1 ) and (512) at 2θ~30.8° and 65.4°, respectively. However, evolution of several additional peaks due to monoclinic WO3 clearly indicate the formation of secondary WO3 phase in the GWO films upon annealing at 900 oC. 21 ACS Paragon Plus Environment


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Surface morphology and topography data obtained from SEM and AFM are presented in Figs. 11 and 12, respectively. In complete agreement with the GIXRD results, SEM and AFM data of the films demonstrate the granular morphology. The topographic images indicate that the GWO surfaces exhibit grain formation on the films. It is also observed that the pore-formation takes place on the films (Figs. 11,12). SEM and AFM data agree with each other in terms of grain size and pore formation; however, the average grain size changes with W-content. Grain size of 38 nm is noted for intrinsic Ga2O3 films while increasing W-content first decreases the grain size and then increases rapidly at higher W-content. The grain size decreases to 32 nm and then rapidly increases to 46 nm for the GWO-100 films, which contain the highest W-content. The changes in surface morphology can be better appreciated in GWO films, where the change from amorphous to granular surface was observed; nevertheless, the pore formation was presented more or less in all the samples after annealing. Intrinsic Ga2O3 films also exhibit the presence of pore formation when annealed at 900 °C; however, the pore formation and coverage seems to increase with increasing sputtering power and, hence, W-content. These observations clearly indicate that the W incorporated into Ga2O3 is responsible for such pore formation upon thermal annealing while the films tend to become fully crystalline due to thermal energy acquired by the annealing process. The data also indicate that the W-content (at%) incorporated into Ga2O3 also plays an important role in the porosity formation upon thermal treatment. As noted in RBS and XPS measurements, the W-content decreases from 12.5% to 11.9%. Thus, the porosity formation effect is dominant after the W-content is decreased from 12.5% to 11.9% upon thermal treatment. This decrease is caused by the diffusion in both inward and outward


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direction by the tungsten in the film. Such diffusion allows the formation of pores of variable size, where the maximum pore size noted was ~27 nm. Similarly, the most dominant morphological change noted is in the samples deposited at a lower sputtering power of 50 W, where the W-content is least. In other words, the lower W-content is responsible for higher porosity density as noted in SEM and AFM data upon thermal annealing of the GWO films. DISCUSSION The results obtained can be conveniently used to discuss the effect of Wincorporation on the crystal structure, surface/interface chemistry, electronic structure and thermal stability of GWO films. First, the observed growth behavior, crystal structure, grain-size variation, and surface morphology evolution in nanocrystalline Ga2O3 as a function of W-content can be explained based on the combined effect of thermodynamic parameters and W-induced structural effects. Note that the growth temperature, which is an important thermodynamic parameter, plays an important role, besides the reactive oxygen partial pressure, in deciding the structural chemistry as well as properties of materials resulting from vapor-transport deposition. If the deposition temperature is low so that the period of the atomic jump process of ad-atoms on the substrate surface is very large, then the condensed species may stay stuck to the regions where they are landing thus leading to an amorphous or amorphous medium.26 The small size grains, spherical in shape, observed in SEM images coupled with the presence of well-resolved peaks indicates the structural order resulting in the formation of nanocrystalline Ga2O3 films at Tc=500 oC. This is due to increased diffusion of ad-atoms leading to a larger rate of atoms joining together and, hence, formation of nanocrystalline films as a result of an increase



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in deposition temperature. The optimum value of Tc required to realize the nanocrystalline films is dependent on the specific oxide material in question. For instance, the onset of crystallization for ZrO2 or HfO2 nanocrystalline films occurs at a temperature of Tc~300 oC58,59 while other oxides such as WO3 or MoO3 can result in the formation of nanocrystalline films even at much lower temperatures, as low as Tc~200 oC.60,61 Thus, the observed higher temperature to deposit nanocrystalline Ga2O3 films by the sputter deposition can be attributed to the difference in material’s physical, chemical and electronic characteristics when compared to other oxides. We now focus our attention to evaluating the structural and morphological behavior of intrinsic and GWO films in order to account for the effect of W-incorporation into Ga2O3 films. The fact that the intensity of XRD peaks diminishing or becoming diffuse indicates the W-incorporation induced amorphization in GWO films. Thus, required temperature of 500 oC to promote nanocrystalline films formation in the case of intrinsic Ga2O3 films is no more valid or sufficient for the case of W-incorporated films. Furthermore, the disordering effect becomes pronounced with a further increase in Wcontent in GWO. However, interestingly as noted in regular XRD measurements, the overall bulk crystallinity of the GWO samples was preserved. Therefore, the most remarkable effect of W-incorporation into Ga2O3 is the tendency to inhibit the crystallization leading to the observed amorphous nature. However, as is evident from GIXRD data, the effect is more pronounced on the surface layers while the bulk of GWO matrix still retains the intrinsic crystal structure of Ga2O3. No other peaks due to any of the W-oxides were detected in XRD patterns. As such, there was no evidence of WO3 or WOx phase formation in the GWO films which allow us to conclude that the W ions at


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the lower content may take the substitutional positions in parent Ga2O3 structure while higher content initiate disordering. Such metal-incorporation induced amorphization was also noted in the case of Ti-doping into WO3.60,62 However, Ti-induced amorphization of the WTiO films was all along the film or complete without any restriction to the surface layers as noted for GWO films in this work. The only difference was that a higher amount of Ti was needed to induce complete amorphization.62 Thus, while Wincorporation maintains the overall crystalline nature of the bulk of the samples, the surface layers become amorphous as is evident in GIXRD patterns, where the patterns become diffuse with increasing W-content. Note that, for a set of carefully designed conditions of incidence angle, the X-ray depth of penetration can be restricted to only a few tens of Angstroms as previously demonstrated for Ga2O3. We relied on the intrinsic gallium oxide information for choosing the experimental conditions in the present work. This is due to the fact that doping of W into Ga2O3 induces a change in the overall density of the films compared to those density values of intrinsic Ga2O3. Nevertheless, by utilizing an incident angle of ≥1o, the x-ray depth penetration can be restricted to a few tens of Angstroms of the sample surface as demonstrated in the previous studies.63 Based on the structural analysis, we formulate the following simple model to account for the observed growth behavior. According to the crystal growth kinetics and the texture evolution of a growing film, the surface energy is the driving factor for the texture evolution.64,65 Thus, the structural configuration and texturing preference in the bulk portion of the Ga-oxide films is fully influenced by the oxide deposit interaction with the Si substrate. However, the surface and strain energy minimization requirements of the films after certain thickness will reduce such interaction, allowing crystal orientation to



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shift as the film continues to grow in thickness. Such phenomena and texturing differences were clearly observed in various oxide films deposited by physical vapor deposition methods, mainly sputtering or laser ablation deposition.26,28,41,42,50,53,55–62 Thus, we believe that, W-incorporation induces a change in the surface and strain energy of the growing GWO film so that the most part of the film underlying is crystalline while the surface layers become amorphous as probed in GIXRD. Further evidence for the proposed mechanism can be derived from the SEM images. The characteristic granular morphology noted for intrinsic Ga2O3 films becomes featureless and smooth with increasing W-content to form GWO films. Since RBS data clearly provide the evidence for W-incorporation uniformly all along the depth of the films, the microscopic imaging analysis is a strong evidence that the W-incorporation effect of amorphization and smooth morphology is more dominant at the surfaces of the GWO films. However, the XPS results provide a clue that the origin of such energy difference and amorphization is mainly due to multiple valence states of W in the GWO films. It was found that the XPS peak broadening was a mixture, where there is a contribution from metallic tungsten (W0). Such species may interfere with growth of ceramic films leading to observed amorphization on the surface layers since depth profiling measurements also provide evidence for some unreacted, metallic tungsten on the film surfaces. In fact, such metallic W unreacted is the content that is responsible to form hydroxyl bonds as seen in XPS data. Finally, the proposed mechanism has been experimentally verified by increasing the deposition temperature and observing the crystal structure of the GWO films. The results indicated that, while Tc=500 oC in the case of intrinsic Ga2O3 films for crystallization, the temperature required was elevated to 800 °C to realize fully crystalline GWO films.


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The chemical analyses made using RBS and XPS agree with each other and indicate that the W-concentration increases with increasing sputtering power to W-target. Additionally, the samples are chemically homogeneous along the depth of the GWO-Si interface to the surface as evidenced in RBS. However, chemical state identification of Ga and W ions in the GWO matrix can only be relied on XPS measurements. Characteristics of Ga 3d core level XPS data, namely the doublet components and their respective energy position, and no changes in either peak shape or BE position, clearly accounts for Ga-O bonds, where Ga ions exist in their highest oxidation state (Ga3+).44–46 Thus, based on XPS analyses, we can conclude that W-incorporation within the 0-12 at% range does not induce any changes in the chemical state of Ga ions. This is quite important to achieve since it has been reported that a large oxygen deficiency in Ga oxide films leads to the formation of a composite system, which exhibits a characteristic insulator to metal transition.66 Furthermore, nonstoichiometric gallium oxide films with very high level of oxygen deficiency can exhibit distinct transport properties (coupled with a superconducting transition at low temperatures), which were accounted only by the presence of Ga metallic clusters in an amorphous Ga2O3 matrix.67 In the present work, absence of lower BE components and further XPS analyses of GWO films clearly rule out of the formation of Ga0 states (metallic Ga). However, chemistry of W in the GWO is quite interesting and not very straight forward. While W 4f7/2 peak at 35.7 eV is definitely a characteristic of W6+ state, unlike Ga 3d core level, the W 4f experiences a peak broadening effect although there is no appreciable BE shift. The chemical state of W is, thus, characterized by a mixture of multiple valence states, which also include some metallic tungsten. Unfortunately, RBS do not shed any light on the W chemical state but



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clearly indicate that W is incorporated all along the film depth on the Si substrate. Based on combined RBS and XPS analyses, we can infer that the W-incorporation is definitely active from early stages to the final stages of the film growth so that W-content is there all in the bulk of the GWO film but some exists in highest oxidation state (W6+) while lower valence states also contribute to the GWO films. The mixed nature of W ion chemical states may form donor levels within the bandgap and account for the significant red-shift in the absorption edge as reported previously.36 Finally, based on the results, we propose a three stage, independent processes, which account for the thermal stability and annealing induced changes in the GWO nanocrystalline films. These are: (a) pore formation leading to porosity in the GWO films, (b) W-diffusion across the interface and (c) the reduction of W atomic content. These phenomena, which occur simultaneously under thermal treatment in GWO films, can be explained as follows. The W-diffusion occurs as per the defect equilibria in the system under the effect of W-incorporation. As gathered from XPS results, W+6 and Ga+3 were the dominant ionic states of W and Ga ions, respectively, in the GWO films. The ionic size of W+6 and Ga+3 are 0.060 nm and 0.062 nm, respectively.68 These values are very close and comparable to each other. Additionally, while Ga2O3 and WO3 both exhibit polymorphism, monoclinic phase is thermodynamically preferred phase for these oxides. This allows, under W-doping up to certain concentration, W ions to act as substitutional defects in the host matrix of Ga2O3. The defect chemistry of the films can be presented in the following equation using Kröger-Vink notation: [1]


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The eq. [1] shows that the charge equilibria for W-incorporation, where W is assumed to be acting as a substitutional defect, is achieved by introducing Ga3+ vacancies on the Ga2O3 matrix.15,37 This process open pathways for W ions to flow from the surface towards the bulk portion of the films, as evident from the results of XPS and RBS measurements. The mechanism for porosity formation and the loss of W-atomic percentage inside the GWO film are explained through the formation of metal-hydroxyl bonds at the surface as previously reported.15 It is important to note that the XPS data revealed the presence of hydroxyl bond formation, which is mainly due to some of the W species forming W-O-H bonding on the surfaces. In order to derive a comprehensive understanding and physico-chemical mechanism involved, the high-resolution XPS data was analyzed using peak pitting and deconvolution of peaks to respective components. Specifically, the O 1s core-level has been analyzed in detail to derive quantification of various components and functional groups. The peaks representing various components are shown in Fig. 13 for as-deposited and annealed GWO samples with variable Wcontent. The de-convoluted peak located at ~532 eV is attributed to the O-H groups, while the peak at ~531 eV corresponds to the oxygen coordinated with W-ions. The quantitative results obtained are presented in Table III. It is clear (Table III) that the annealing process induces a higher level of hydroxyl interaction with GWO sample surface. This is attributed to the moisture adsorbed by thermally induced pore formation in the GWO films under annealing. Similarly, a shift in the W6+ contribution is evident after the films are annealed (Fig. 13b). This is due to the evaporation of WO3 molecules from film surface which in turn is accounts for the reduction in the overall W-content in



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the GWO films as noted in chemical analyses. Under thermal annealing, tungsten oxides become highly volatile under the presence of water vapor;69 the evaporation of W-O-H molecules on the film surface diminishes the overall W-content, while contributing to the pore generation in the films. GIXRD results demonstrated the presence of WO3 formation on the surface after the films were exposed to high temperature annealing; these secondary phase formation of tungsten oxide was expected due to difference in electronegativity of both Ga and W ions. Pore formation generally increases the surface to volume ratio which is imperative in certain physical-chemical applications such as catalysis and chemical sensing. Note that both WO3 and Ga2O3 were considered as feasible materials for photocatalytic applications,22,24,70,71 where the functionality can be significantly enhanced by the increasing the surface to volume ratio which in turn enhances the surface chemical reactivity72. More work needed to be performed in order to derive a more detailed account of a composition-property correlation for the case of W-incorporated gallium oxide and quantification of improvement of such properties and device performance. However, the present structural-chemical analysis performed clearly implies possible advantages of GWO films for the aforementioned applications. This can be further exemplified by the available recent literature, where several authors emphasized the importance of the structural-chemical interaction of metallic ions doped into Ga2O3 for tuning the physico-chemical properties. For instance, Shimura and Yoshida presented the influence and key role of a secondary phase formation during Zn doping into gallium oxide on the photocalatylic activity of the material.22 Similarly, Lopez et. al. showed the importance of In-doping on the structural-chemical characteristics of Ga2O3


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nanostructures, suggesting the importance of metallic doping into the growth mechanism and enhanced photo luminescence of the material.9 Furthermore, is appears from XPS analyses that the hydrogen adsorption capacity increases whenever the W-doping is presented. While further analysis is needed in that direction, the level of adsorption increases due to the presence of hydrophilic WO3. To corroborate the presence of the two mechanisms discussed in this work (W-diffusion and W-content loss), XPS depth profiling measurements were performed on a representative sample (GWO-100) before and after high temperature treatment. The results are shown in Fig. 14. It is evident that change in the overall atomic concentration of W decreased after annealing which can be explained by the volatilization of W-hydroxyl bonds. Furthermore, it can be noticed that the oxygen content in the film decreases associated with an increase in the thickness of the Si-GWO interface after annealing. The change of oxygen content is related to the Wcontent loss, which can be explained through the change of the oxygen to metal ratio decreasing from 3 for WO3 to 1.5 for Ga2O3 which is the predominant oxide after the films are annealed. In the case of the interface broadening, it is attributed to the diffusion of Si from the substrate towards the interfacial region of the film. Finally, based on the observed XPS and RBS data of the GWO films upon annealing, the chemical changes in terms of W-content in the GWO films are schematically represented in Fig. 15. The diagram presents the W-content variation along the depth of the film. Also, the data is presented to provide a comparison between as-deposited and annealed GWO films. The changes in overall W-concentration after the annealing process and the regions where the GWO/Si intermixing occurs are evident. Furthermore, preferred location of tungsten in



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the film after the heat treatment is indicated by the parabolic behavior that the Wconcentration curve follows. CONCLUSIONS An approach is presented to incorporate tungsten into Ga2O3 with a control over phase and structural chemistry by carefully controlling the processing conditions in cosputter deposition of respective targets. Varying the sputtering power to the W-target from 0 to 100 W produces GWO films with W-concentration varying from 0 to 12 at%. Crystal structure, surface and interface chemistry, valence state of Ga and W, morphology and chemistry of intermixing were evaluated for GWO films as a function of W-content. Intrinsic Ga2O3 films were stoichiometric, nanocrystalline with a β-phase monoclinic structure. GWO films with varying W from 0 to 12 at% also crystallize in monoclinic β-Ga2O3 phase suggesting that W-atoms may be occupying substitutional positions without any perturbation to the crystal structure. This feature is attributed mainly to similarity of ionic radii of W and Ga ions and adoptability of monoclinic structure by their respective oxides in their highest oxidation state. However, Wincorporation induces surface-specific amorphization, which could be due to surface energy difference of the GWO system compared to intrinsic Ga2O3. The XPS results provide a clue that the origin of such energy difference and amorphization is mainly due to multiple valence states of W in the GWO films. RBS analyses confirm the uniform distribution of W-content in the GWO films. GWO films were stable up to a temperature on the order of 900 oC, at which point W-interdiffusion activates. However, the interdiffusion is significant only in GWO samples with highest W-content (12 at%) while samples with lower W-content exhibit thermo-chemical stability. Thermal stability of


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GWO films evaluated at 900 oC indicates thermal-annealing induced secondary phase (WO3) formation. A transition to mesoporous structure coupled with W interdiffusion occurs due to thermal annealing as derived from the chemical analyses at the GWO films’ surface as well as depth-profiling towards the GWO-Si interface. Thermally induced morphological changes are dependent on W-concentration in the GWO films. Thermallyinduced W-diffusion and intermixing also induces pore formation with variable size. The maximum pore size noted was ~27 nm for GWO films with highest W-content. XPS analyses indicate the electronic charge redistribution is dominated by the hydroxyl groups and W-chemistry upon thermal annealing. Combined use of RBS and XPS indicate that the extent of diffusion and intermixing layer depth is dependent on W-content in the GWO films; pronounced thermally induced W-diffusion and with Si-W-Ga2O3 intermixing is noted only in GWO samples with highest (12 at%) W-incorporation. A model has been formulated to account for the mechanism of W-incorporation, thermal stability and interdiffusion via pore formation in GWO films. The structure and morphology of the GWO films exhibit a transition from amorphous to fully crystalline state with a high content of pores in the case of W-containing films. Corroborating with RBS results, these changes were, however, remarkable only in samples with high Wcontent, whereas intrinsic Ga2O3 did not show any significant changes due to the annealing process. Conglomeration of present work and available data in the literature suggest the functional relationship between W-incorporation, surface and interface chemistry, thermally induced effects as derived in this work will have implications in tuning the physio-chemical properties of GWO films for a given technological application. The porous structure formation and associate pore-structure-property



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quantification requires further attention to further tune the morphology and utilize the benefits of morphology-controlled GWO materials for future applications. Acknowledgements

Authors at the University of Texas at El Paso acknowledge, with pleasure, support from the National Science Foundation (NSF) with grant # ECCS-1509653. Part of the work and sample analysis was also performed by support from the National Science Foundation (NSF) with NSF-PREM grant # DMR-1205302. A portion of the research was performed using Environmental Molecular Sciences Laboratory (EMSL), a DOE Office of Science User Facility sponsored by the Office of Biological and Environmental Research and located at Pacific Northwest National Laboratory. XPS data were obtained at the Materials Research Laboratory (MRL) Shared Experimental Facilities at UCSB supported by the MRSEC Program of the NSF under Award No. DMR 1121053; a member of the NSF-funded Materials Research Facilities Network.


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Figure Captions Figure 1

RBS data of GWO films deposited with variable W-content. Data shown include experimental curves (circles) and SIMRA simulated curves (lines) for GWO films. The backscattered ions due to W, Ga, O and Si and their respective energy position are as indicated. A reasonable agreement between the experimental and simulated curves can be seen.

Figure 2

XPS survey spectra of GWO films with variable W-concentration. The peaks due to various elements and their respective energy positions as indicated.

Figure 3

High resolution XPS data of Ga, W and O core level peaks in GWO films deposited with variable W-concentration. (a) Ga 3d high resolution XPS spectra, (b) W 4f core-level XPS spectra and (c) O 1s high resolution XPS spectra.

Figure 4

GIXRD spectra GWO films deposited under variable sputtering power to the W-target. The peaks become less intense and diminish with increasing sputtering power to the W-target i.e., increasing Wconcentration.

Figure 5

SEM micrograph of GWO films. The featureless, characteristic smooth surface morphology evolution with increasing W-concentration is evident.

Figure 6

RBS data of GWO films annealed at 900 oC. Data shown include experimental curves (circles) and SIMRA simulated curves (lines) for GWO films. The backscattered ions due to W, Ga, O and Si and their



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respective energy position are as indicated. A reasonable agreement between the experimental and simulated curves can be seen. Figure 7

Schematic representation of thermally induced W-diffusion and depth penetration into the Si substrate with Si-W-Ga2O3 intermixing at the interface. The GWO-Si interface before annealing (a) can be compared and contrasted with the W-concentration gradient due to intermixing upon annealing in GWO-100 sample (b).

Figure 8

XPS survey spectra of GWO films after annealing at 900° C. The data shown are for variable W-concentration.

Figure 9

High resolution XPS data of Ga, W and O core level peaks in GWO films annealed at 900 oC. (a) Ga 3d high resolution XPS spectra, (b) W 4f core-level XPS spectra and (c) O 1s high resolution XPS spectra

Figure 10

GIXRD spectra GWO films annealed at 900 oC. It is evident that the peaks corresponding to β-Ga2O3 appear. Peaks due to WO3 are also noted as indicated due to the secondary phase formation in GWO films upon thermal annealing at 900 oC.

Figure 11

SEM micrographs of GWO films after annealing. The morphology changes upon annealing are clearly seen in the micrographs.

Figure 12

AFM micrographs of GWO films after annealing. The morphology changes upon annealing corroborate with SEM data shown in Fig. 11.

Figure 13

XPS peak fitting and deconvolution of oxygen 1s core level. The data are presented for GWO samples with variable W-content. The data shown are for (a) as-deposited and (b) after annealing at 900 °C.


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Quantitative information derived from analysis is presented in Table III. Figure 14

XPS depth profile of representative GWO films. The data shown are for (a) as-deposited and (b) after annealing at 900 °C for 1 hour. The data obtained for Ga, W, O and Si are shown. The error of estimation is ±0.01 at% for metals (Ga, W, and Si) and ±0.1 at% for oxygen.

Figure 15

Schematic diagram of the proposed mechanism and W-concentration profile as evidenced by the XPS and RBS data. A comparison of Wconcentration in the GWO samples before and after annealing at 900 oC is presented.



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Figure 1


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Figure 2

600

500

400

300

200

B.E. (eV)


100

W4f Ga3d

Ga3s

W4d5/2

C1s

W4d3/2

Ga LMM

Ga3p

GWO-100 GWO-50

O1s

CPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0


Figure 3

1: Ga 3d5/2 2: Ga 3d3/2

Ga 3d

2

1

O 2s CPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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GWO-100

O 2s

1 2

(a)

GWO-50

26

24

22

20

B.E. (eV)


18

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(b)

W4f7/2

(c)

O1s

W4f5/2

CPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


GWO100 W5p

GWO100 GWO-50 GWO-50 44

40

36

534

B.E. (e.V)


531

528


Figure 4

GWO-50

Intensity (a.u)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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GWO-75

GWO-100 (512)

(40-1) (111)

25

30

35

GWO-0 (Ga2O3)

40

45

50

2θ (Degrees)


55

60

65

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Figure 5



Figure 6 Energy (keV) 400 7000

800

1200

Energy (keV) 1600

2000

400 7000

experimental simulated

6000

+

4000

Ga Si O

4000 3000

2000

2000

1000

1000

400

Si

600

800

0 200

1000

Ga

400

Channel Energy (keV) 400 7000

800

1200

1600

2000

400 7000

800

1000

+

800

1200

1600

2000


6000

W

+

2MeV He → Ga2O3//Si (100) 5000

600

W

Channel Energy (keV)


6000

2000

GWO-50 Annealed

O

0 200

1600

+

5000

Counts

Counts

1200

2MeV He → Ga2O3//Si (100)

GWO-0 (Ga2O3) Annealed

3000

800 experimental simulated

6000

2MeV He → Ga2O3//Si (100) 5000

2MeV He → Ga2O3//Si (100)

GWO-75 Annealed

5000 GWO-100 Annealed 4000

Si

3000

Counts

4000

Counts

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W

O Ga

2000

Si

3000

O 2000

Ga 1000 0 200

1000

400

600

800

1000

0 200

400

600

Channel

Channel


800

1000

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Figure 7



Figure 8

GWO-50

CPS

GWO-75

O1s

600

500

400

300 B.E. (eV)


200

100

W4f Ga3d

Ga3p

Ga3s

W4d5/2

100W-W

Ga LMM C1s W4d3/2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Figure 9

Ga 3d

O1s

W4f7/2

(a)

(b)

(c)

W4f5/2

CPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


GWO100

GWO-75

GWO-50

24

21

18

42

39

36

B.E. (eV)


33

534

531

528


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 10


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Figure 11



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 12


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Figure 13

(a) Core Spectra 6+ W (O1s) O-H (O1s) 3+ Ga (O1s)

50W-W Counts (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


As Deposited GWO Films

100W-W

536

534

532

530

B.E. (eV)


528


Annealed GWO Films

(b) 50W-W

Counts (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 62 of 68

Core Spectra O-H (O1s) 6+ W (O1s) 3+ Ga (O1s) 75W-W

100W-W

536

534

532 B.E. (eV)

530


528

Page 63 of 68

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 14



6

W at% annealed W at% as-grown

W-O-H evaporation

9

GWO/Si Interface

Figure 15

Concentration (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 64 of 68

6+ Vga induced by W

W-diffusion Region of Stable W- (at%) Concentration

3 W-diffusion

0

10

20

30

Depth (nm)


40

50

Page 65 of 68

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table I: Chemical composition and thickness of GWO films as determined from RBS data. The error of these estimated quantities are as indicated. Sputtering Power (W)

Sample ID

W (±0.02 at%)

0

GWO-0

0.0

Thickne ss (±1 nm) 32

50

GWO-50

8.4

39

75

GWO-75

9.5

42

100

GWO-100

12.5

53



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 66 of 68

Table II: Tungsten concentration in GWO films after annealing at 900 oC. Data obtained from RBS and XPS are presented and compared. W-Sputtering Power (Watts)

Sample ID

W- RBS (±0.02 at%)

W-XPS (±0.01 at%)

0

GWO-0

0

0

50

GWO-50

4.3

1.0

75

GWO-75

6.4

1.7

100

GWO-100

11.9

2.0


Page 67 of 68

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table III: Quantitative information of various functional groups as obtained from XPS peak fitting and analysis. The data indicates the functional groups contributing to the oxygen peak as obtained from O 1s peak fitting. The error of estimation is ±0.1% for the functional groups.

Samples As-deposited (Pw) 50 W 100 W Annealed (Pw) 50 W 75 W 100 W

O1s O-H (%)

W6+ (%) Ga3+ (%)

9.6 7.0

38.1 31.5

52.3 61.6

17.2 25.1 22.4

24.6 23.1 20.2

58.2 51.8 57.5



6

GWO/Si Interface

9

W-O-H evaporation

TOC Graphic

Concentration (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

900 oC

6+ Vga induced by W

Region of Stable W- (at%) Concentration

3 W-diffusion

0

10

W-diffusion

Wat% annealed Wat% as-grown

20

30

40

50

Depth (nm)

ACS Paragon Plus Environment

Page 68 of 68

Tungsten Incorporation into Gallium Oxide - American Chemical Society

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