β-Ga2O3 has been considered as a promising semiconductor material for next-generation high power electronic devices due to its ultra-wide bandgap energy (4.8-4.9 eV), theoretically predicted high breakdown field strength (8 MV/cm), and for the ease of mass production of free-standing single crystal native substrates [1,2]. The demonstration of n-type doping in β-Ga2O3 over a wide range of electron concentrations (10 16 -10 20 cm -3 ) with decent transport properties such as the mobility close to 200 cm 2 /V.s at room temperature and >10 4 cm 2 /V•s at cryo-temperature with ultra-low background concentration (10 14 ) reveal this material's immense potential in high power electronic and deep ultraviolet optoelectronic applications [3][4][5][6][7][8][9][10][11][12][13][14][15][16][17][18][19]. Another key advantage of β-Ga2O3 is its capability of the bandgap engineering by alloying with Al2O3 which can extend the bandgap energy up to 7.24 eV in monoclinic phase [20]. Due to the tunability of bandgap energy over a wide range, β-(AlxGa1-x)2O3 alloys possess a great promise for next generation high power and radio frequency electronic applications.
Considering the enormous potential of β-(AlxGa1-x)2O3 alloys in electronic and optoelectronic applications, several studies have been conducted on the development of β-(AlxGa1-x)2O3 epitaxy [21][22][23][24][25][26][27][28][29][30][31][32][33][34] and β-(AlxGa1-x)2O3/β-Ga2O3 heterostructure-based lateral devices [35][36][37][38][39] grown by different growth techniques including molecular beam epitaxy (MBE) [21][22][23] and metalorganic chemical vapor deposition (MOCVD) [24][25][26][27][28][29][30][31][32][33][34]. MOCVD grown phase pure β-(AlxGa1-x)2O3 thin films on (010) (x ≤ 35%) [24,25,29,33], (100) (x ≤ 52%) [26,27] and (2 � 01) (x ≤ 48%) [28,33] oriented β-Ga2O3 substrates have been demonstrated with a great control of elemental compositions, uniformity, and high purity. β-(AlxGa1-x)2O3 alloy with < 80% Al composition is predicted to have a breakdown field strength up to 16 MV/cm [40], which is significantly higher than the fields achievable from SiC or GaN based devices, even with nominally lower content of Al. Moreover, high-Al composition in phase pure β-(AlxGa1-x)2O3 layer is advantageous as it offers opportunity to generate two-dimensional electron gas (2DEG) by modulation doping at the β-(AlxGa1-x)2O3/β-Ga2O3 interface owing to its tunable and large conduction band offset [35][36][37][38][39].
Excellent transport properties have already been demonstrated for modulation doped field effect transistors (MODFETs) based on β-(AlxGa1-x)2O3/β-Ga2O3 heterostructures. A room temperature mobility of < 180 cm 2 /Vs and a cryo-temperature peak mobility of < 2790 cm 2 /Vs were measured at a 2DEG sheet charge density of < 5 × 10 12 cm -2 for Al compositions x ≤ 0.20 [35,36]. The successful operation of these MODFET devices shows great potential for developing high performance transistors using β-(AlxGa1-x)2O3/β-Ga2O3 (x ≤ 0.20) heterostructures.
While several studies on the epitaxial growth of β-(AlxGa1-x)2O3 thin films over a wide range of Al compositions have been reported, the investigation of the n-type doping of β-(AlxGa1-x)2O3 is still limited. Recently, theoretical studies based on hybrid density functional theory calculations have predicted Si as the most efficient shallow donor for high Al-content β-(AlxGa1-x)2O3 alloys, effective up to 70-85% Al compositions [41,42]. Our previous experimental study on MOCVD grown Si doped β-(AlxGa1-x)2O3 thin films grown with 6.3-33.4% Al compositions exhibited promising room temperature (RT) mobility (42-108 cm 2 /Vs) at doping concentrations ranging between 10 17 -10 18 cm -3 [24]. Another investigation on the uniformly Si doped MOCVD β-(Al0.26Ga0.74)2O3 films showed higher carrier concentrations of 6×10 18 -7.3×10 19 cm -3 with corresponding RT mobilities of 53-27 cm 2 /V.s [43]. Although these studies indicated the possibility for n-type Si doping in β-(AlxGa1-x)2O3 thin films for a wide range of doping concentrations and Al compositions, the measured transport results can be significantly impacted by the formation of a modulation doped channel in β-Ga2O3 and by the electron accumulation at
The β-(AlxGa1-x)2O3 films with 6%, 11% and 18% Al compositions were grown on Fe doped semi-insulating (010) β-Ga2O3 substrates by using Agnitron Technology Agilis R&D MOCVD system. The β-Ga2O3 substrates were purchased from Novel Crystal Technology, Inc. The Si doped ~350-450 nm thick β-(AlxGa1-x)2O3 films were grown on ~65 nm thick Mg-doped semiinsulating β-Ga2O3 buffer layer on top of (010) β-Ga2O3 substrates. The Mg doped β-Ga2O3 buffer layer was grown at a growth temperature of 700 °C with a chamber pressure of 60 torr. The growth temperature of 880 °C and the chamber pressure of 20 torr were used to grow the β-(AlxGa1-x)2O3 layers on top of Mg doped β-Ga2O3 buffer layer. Triethylgallium (TEGa), Trimethylaluminum (TMAl), Bis-(cyclopentadienyl)-magnesium (Cp2Mg) and pure O2 were used as Ga, Al, Mg and O precursors, respectively. Ar was used as the carrier gas. The Mg/Ga molar flow rate ratio was 6.69 x 10 -3 . Al composition of 6%-18% was obtained by varying the [TMAl]/[TEGa+TMAl] molar flow rate ratio from 1.08% to 3.71%. The O2 flow rate was set at 500 sccm for all growths. Silicon dopants were introduced into the chamber by flowing diluted silane (SiH4). Different doping concentrations of β-(AlxGa1-x)2O3 thin films were obtained by controlling the mass flow controllers in the silane and dilution lines. The silane flow rate was tuned from 0.53 to 100.1 nmol/min. All the substrates were ex-situ cleaned by solvent before loading to the MOCVD reactor.
Any potential contamination on the substrate surface was removed by performing a high temperature in-situ annealing at 920 °C for 5 mins duration under O2 atmosphere prior the growth.
The Al composition and crystalline structure of the β-(AlxGa1-x)2O3 films were evaluated by high resolution XRD by using a Bruker D8 Discover (Cu Kα source, λ=1.5418 Å). Surface morphologies and roughness were characterized by optical (Huvitz HRM-300) and atomic force (AFM, Bruker Icon 3) microscopy. The β-(AlxGa1-x)2O3 film thicknesses were estimated from the cross-sectional field emission scanning electron microscope (FESEM) images of the reference (AlxGa1-x)2O3 samples grown on the co-loaded c-plane sapphire substrates and by using high resolution STEM imaging. A Thermo Fisher Scientific Themis-Z scanning transmission electron microscope (operated at 200 kV) was used for high resolution STEM imaging. The elemental concentrations of β-(AlxGa1-x)2O3 were characterized by SIMS. An HMS 3000 Fast Hall measurement system (Ecopia) was used to measure the carrier concentration and electron Hall mobility at room temperature. 20/100 nm Ti/Au contacts were deposited on four corners of the sample to create the van der Pauw setup for Hall measurement. Temperature dependent Hall measurement was performed using two custom built systems. For the measurement below room temperature, an electromagnet with a vacuum cryostat with closed cycle helium (He) refrigerator was used. An electromagnet with a quartz tube and silicon carbide heater was used for the measurement above room temperature. Substitutional carbon on a cation/oxygen site acts as a shallow donor in Ga2O3 as its positive charge states are stable over the entire Ga2O3 bandgap which was predicted by theoretical study [44].
However, experimentally in our MOCVD grown high quality β-Ga2O3 homoepitaxial films [6],
the Si concentration extracted from quantitative SIMS characterization matched well with the net Nd-Na concentration, regardless of the existence of C background level, implying that C may not act as an active shallow donor in β-Ga2O3. While the growth conditions may also influence the role of C incorporated in β-Ga2O3, further investigation is required to confirm about the exact state of C in β-Ga2O3 [45,46]. Nevertheless, both C and H turn into compensating acceptors in β-(AlxGa1-x)2O3 as the band gap widens due to higher Al incorporation, predicted by recent DFT calculations [42]. As compared to configurations such as the carbon occupying the octahedral cation sites or the oxygen sites, the C on the tetrahedral Ga site is found to be most stable in the positive charge state within Ga2O3 bandgap from DFT studies, deriving that C act as a shallow donor in Ga2O3. Similar to Ga2O3, the theoretical DFT calculations also predicted that carbon impurity prefers to occupy the substitutional tetrahedral cation site in monoclinic θ-Al2O3. In case of neutral charge state of CAl(I) (C on tetrahedral site) in θ-Al2O3, carbon moves beyond the plane of its oxygen neighbors towards a next-nearest Al(I) neighbor, causing a reduction of the C-Al distance from 3.5 to 2.6 Å. As a result, the charge of the localized state is confined between the carbon and that next-nearest-neighbor Al atom, indicating that CAl(I) (with a corresponding Kohn-Sham state lying 4.42 eV below the conduction band minimum) is not a shallow donor in θ-Al2O3.
By interpolating the (+/-) charge-state transition level of the C(I) negative-U center (carbon substituting the tetrahedral Ga or Al site) between β-Ga2O3 and θ-Al2O3, carbon is found to act as acceptor with only 5% Al incorporation. On the other hand, carbon replacing the octahedral Ga or Al site (C(II)) starts acting as an acceptor at > 51% Al compositions. However, carbon on oxygen sites (CO) are predicted to act as compensating acceptor for both Ga2O3 and Al2O3 under n-type doping conditions. Another common impurity in MOCVD growth system-hydrogen also acts as compensating acceptor in θ-Al2O3 as they are found to be stabilized in negative charge states, while all configurations of H substituting on the O site or interstitial site are found to act as shallow donor in β-Ga2O3. In addition, recent DFT calculation also predicted that the Al energetically prefers to occupy the octahedral Ga2 site and is forced to occupy the tetrahedral Ga1 site when the Al composition in β-(AlxGa1-x)2O3 increases [30]. However, due to the surface reconstructions and kinetic limitations during the epitaxial growth, Al atoms are observed to occupy both octahedral and tetrahedral sites in β-(AlxGa1-x)2O3 alloy, eventually leading to a structure with higher energy by the increase of total number of Al atoms substituting on the Ga1 site. The lack of control over Al site incorporation results in the formation of extended (planar) defects. The divacancyinterstitial complexes, comprising of one cation interstitial atom paired with two cation vacancies, were predicted to be compensating acceptors in β-Ga2O3 [47]. The higher density of such planar defects in β-(AlxGa1-x)2O3 alloy driven by the local increase of Al atoms might also contribute as compensators in Si doped β-(AlxGa1-x)2O3 thin films. Experimentally, the lower Si incorporation efficiency along with higher level of compensations due to the increase of C and H impurity concentrations and planer defects can potentially contribute to the reduction of effective net doping in β-(AlxGa1-x)2O3 films as evidenced by the comprehensive characterization of room temperature Hall measurements and SIMS. Along with the reduction of carrier concentration, the room temperature Hall mobility also drastically decreases with the increase of Al composition for both cases with different silane flows, which can be attributed to the increase of both alloy scattering and divacancy -cation interstitial defect density developed due to incorporation of higher Al concentration [24].
As the Si doping in β-(AlxGa1-x)2O3 films with higher Al compositions are found to be challenging due to the increase of impurity concentrations, the role of different silane flow rates on transport properties of β-(AlxGa1-x)2O3 films with relatively low Al compositions are investigated. Figure 3(a) shows the room-temperature Hall mobility and resistivity versus carrier concentration for β-(AlxGa1-x)2O3 films grown with 6% Al composition for different silane flow rates ranging between 1.17-14.35 nmole/min. Increasing the silane flow from 1.17 to 14.35 nmole/min leads to the increase of carrier concentration from 6.56×10 17 to 1.54×10 19 cm -3 .
However, the Hall resistivity and the room temperature mobilities are found to decrease with increasing silane flow. As the doping concentration increases from 6.56×10 17 to 1.54×10 19 cm -3 , the increased ionized impurity scattering becomes the dominant factor [3], which limits the roomtemperature mobility from 101 to 54 cm 2 /V.s. Following the room temperature Hall measurements, a selected representative 6% Al composition sample was characterized by temperature-dependent Hall measurement in the van-der-Pauw geometry to probe the donor activation energy, as shown in Figure 3(b) and 3(c). The temperature was varied from 40 to 300 K. The room-temperature Hall carrier mobility was measured as 101 cm 2 /V•s at a carrier concentration of 6.56×10 17 cm -3 . A peak mobility of 1157 cm 2 /V•s at a carrier concentration of 2.30×10 17 cm -3 was measured at 65K. The donor activation energy (ED) of 8 meV with donor (ND) and compensation (NA) concentrations of 1.15×10 18 cm -3 and ~ 4×10 16 cm -3 , respectively, were extracted by fitting the temperature dependent transport data using the charge neutrality equation (𝑛𝑛
), where NA is the concentration of acceptors acting as compensators, ND is the donor concentration, and ED is the donor activation energy. The extracted activation energy of 8 meV for 6% Al composition sample is found to be close to the Si donor energy of β-Ga2O3 with similar doping density [48]. It is worth noting that the increase of Al composition can result in higher Si donor activation energy in β-(AlxGa1-x)2O3 due to the widening of the bandgap, which is clearly evidenced by the increase of donor activation energy from 8 to 36 meV, extracted from 6% and 18% Al composition samples, respectively, as shown later in Figure 4. Apart from the donor activation energy, the compensation concentrations extracted from the fitted data are also observed to increase from ~4×10 16 to ~8×10 16 cm -3 as the Al composition in β-(AlxGa1-x)2O3 films increases from 6% to 18% (Figure 4), which is consistent with the results from SIMS characterization. Previously, the Si donor energy of β-Ga2O3 was found to decease with the increase of carrier density [48]. The Si donor activation energy of 16 meV from a β-(Al0.06Ga0.94)2O3 film grown with a doping concentration of 1.98×10 17 cm -3 was also extracted from the temperature dependent Hall measurement. As the doping concentration of β-(Al0.06Ga0.94)2O3 decreases from 6.56×10 17 cm -3 to 1.98×10 17 cm -3 , the donor activation energy was found to increase from 8 to 16 meV, which is consistent with the findings from β-Ga2O3 films investigated with different carrier concentrations [48]. Similar trend of a sharp reduction of Si activation energy with the increase in Si doping concentration was also observed in highly Si-doped (5.2x10 18 cm -3 < Nd < 1.5x10 19 cm -3 ) AlGaN films grown with 70% Al composition [49].
The growth of Si doped β-(AlxGa1-x)2O3 films on unintentionally doped β-Ga2O3 buffer layer can lead to the formation of a modulation doped channel in β-Ga2O3 or electron accumulation at β-(AlxGa1-x)2O3/β-Ga2O3 interfaces. In addition, the significant peak of Si accumulation at the epilayer/substrate interface, as observed in Figure 2(b), can also add complications in transport characteristics. Such strong Si peak at the substrate growth interface was also observed in previous studies on homoepitaxial β-Ga2O3 thin films [3,6,17]. Owing to its lowest formation energy as compared to other cation site acceptors, Mg is considered as one of the most promising acceptors in β-Ga2O3 with relatively shallow acceptor level [50]. Mg doped β-Ga2O3 buffer layer was found to be highly effective in suppressing the accumulated charges from epi-layer/substrate interface [17]. In this work, the Si doped β-(AlxGa1-x)2O3 films were grown on Mg-doped β-Ga2O3 buffer layer to compensate the charges from the epi-layer/substrate and β-(AlxGa1-x)2O3/β-Ga2O3
interfaces. The role of Mg-doped β-Ga2O3 buffer layer in suppressing the interface charges as compared to the UID β-Ga2O3 layer are probed by the temperature dependent carrier concentration. To investigate the influence of different carrier concentrations of β-(AlxGa1-x)2O3 thin films on the surface morphology, AFM was performed for 6% Al composition samples grown with different doping levels as shown in Figure 5 (a-c). All the films show smooth surface morphologies with RMS roughness ranging between 0.50-1.74 nm. Although similar surface morphologies are observed, the surface roughness is found to increase from 0.50 to 1.74 nm as the doping concentration increases from 6.56×10 17