High-mobility hydrogenated polycrystalline In2O3 (In2O3:H) thin-film transistors

Oxide semiconductors have been extensively studied as active channel layers of thin-film transistors (TFTs) for electronic applications. However, the field-effect mobility (μFE) of oxide TFTs is not sufficiently high to compete with that of low-temperature-processed polycrystalline-Si TFTs (50–100 cm2V−1s−1). Here, we propose a simple process to obtain high-performance TFTs, namely hydrogenated polycrystalline In2O3 (In2O3:H) TFTs grown via the low-temperature solid-phase crystallization (SPC) process. In2O3:H TFTs fabricated at 300 °C exhibit superior switching properties with µFE = 139.2 cm2V−1s−1, a subthreshold swing of 0.19 Vdec−1, and a threshold voltage of 0.2 V. The hydrogen introduced during sputter deposition plays an important role in enlarging the grain size and decreasing the subgap defects in SPC-prepared In2O3:H. The proposed method does not require any additional expensive equipment and/or change in the conventional oxide TFT fabrication process. We believe these SPC-grown In2O3:H TFTs have a great potential for use in future transparent or flexible electronics applications. The field-effect mobility of oxide semiconductor thin-film transistors is not sufficiently high compared to silicon thin-film transistors. Magari et al. use a low-temperature solid-phase crystallization process to fabricate In2O3 thin-film transistors with mobility comparable to silicon counterparts.

W ide-bandgap oxide semiconductors (OSs) have been extensively studied as active channel layers of thinfilm transistors (TFTs) for next-generation flat-panel displays 1,2 , nonvolatile memories 3 , inverters 4 , various sensors 5,6 , Schottky devices 7,8 , and so on. Among OSs, amorphous In-Ga-Zn-O (a-IGZO) TFTs have now become the backplane standard for active-matrix liquid-crystal displays (AMLCDs) and active-matrix organic light-emitting diode (AMOLED) displays because of their reasonable field-effect mobility (μ FE ) of over 10 cm 2 V −1 s −1 , extremely low leakage current, low process temperature (<350°C), and large-area scalability 9,10 . Although the μ FE value of a-IGZO TFTs is more than ten times higher than that of hydrogenated amorphous Si (a-Si:H) TFTs (<1 cm 2 V −1 s −1 ), it is not sufficiently high to compete with that of low-temperature-processed polycrystalline-Si (LTPS) TFTs (50-100 cm 2 V −1 s −1 ) 11 . The main disadvantages of LTPS TFTs are a relatively high process temperature (450-550°C) and an expensive crystallization process. The high μ FE values of OS TFTs mean these devices could be used in fields that have been dominated by LTPS TFTs and in transparent and flexible devices that are incompatible with Si. Numerous types of approaches to enhance the μ FE value of OS TFTs have been investigated, including cation composition 12,13 , multiple channel structures 14,15 , dual-gate architecture 16,17 , metal capping layer structures 18,19 , post treatment 20,21 , and their combination. Among these, cation composition control is the most promising method. It requires no extra complex process for integrating OS TFTs. In-rich OSs have been studied extensively. The large spatial spread of the In 5s orbital with a large overlap can provide a facile electron transport path with a low electron effective mass 22 . However, undoped In 2 O 3 films exhibit a high background electron concentration (10 19 -10 21 cm −3 ) 23,24 , which is attributable to the presence of native defects, such as oxygen vacancies, making it difficult to control the threshold voltage of the TFTs 25,26 . To suppress the carrier concentration in In 2 O 3 , elements such as Ga, Hf, Si, Al, and W were added, which have large bond dissociation energies with oxygen 12 . Many AOS TFTs have been explored with multicomponent oxide semiconductors, such as In-W-Zn-O 27 , Al-In-Sn-Zn-O 28 , and In-Ga-Zn-Sn-O 29 . However, multicomponent oxides complicate the composition control of the deposited film. Moreover, multimetal cations cause potential fluctuation near the conduction band minimum, which might hinder electron transport 30 .
Recently, crystalline OSs have been proposed to enhance the carrier mobility because the disorder-induced subgap states can be suppressed via lattice ordering. Yang et al. reported a μ FE value of 60.7 cm 2 V −1 s −1 for a TFT obtained using polycrystalline In-Ga-O annealed at 700°C 31 . Although high annealing temperatures result in better electrical properties of the oxide active channel layer, such high temperatures are unsuitable for device application on glass or plastic substrates. Our group reported a μ FE value of 50.6 cm 2 V −1 s −1 for a TFT obtained using hydrogenated polycrystalline In-Ga-O formed via solid-phase crystallization (SPC) at 300°C 32 .
This study proposes a simple material and a simple process to obtain high-performance TFTs, namely hydrogenated polycrystalline In 2 O 3 (In 2 O 3 :H) TFTs grown via the low-temperature SPC process. In 2 O 3 :H TFTs fabricated at 300°C exhibit superior switching properties with µ FE = 139.2 cm 2 V −1 s −1 , a subthreshold swing (SS) of 0.19 V dec −1 , and a threshold voltage (V th ) of 0.2 V. The hydrogen introduced during sputter deposition plays an important role in enlarging the grain size and decreasing the subgap defects in SPC-prepared In 2 O 3 :H. The proposed method has great potential for future transparent or flexible electronics applications. . This result indicates that H 2 addition suppresses the growth of crystallites during deposition. After annealing at 250°C in nitrogen for 1 h, the amorphous phase of In 2 O 3 :H changed to the crystalline one with the (222) preferred orientation. The angles of the diffracted peaks are in good agreement with the In 2 O 3 powder data (ICSD code: 14388). Moreover, the crystallized films exhibited smaller fullwidth at half-maximum values of the (222) reflection than the film deposited without hydrogen introduction, indicating larger crystallite sizes and smaller strains in the In 2 O 3 :H films. Figure 1c-h depict the EBSD images along the normal direction for the In 2 O 3 and In 2 O 3 :H films with and without annealing at 250°C in nitrogen. For the as-deposited films (Fig. 1c-e), a randomly oriented small grain structure embedded in the amorphous matrix can be observed in the In 2 O 3 film without H 2 introduction. The grain structure disappeared upon increasing the R[H 2 ] value to 5%. In contrast, a huge grain structure appeared for In 2 O 3 :H deposited at R[H 2 ] = 5% after annealing at 250°C (Fig. 1h), indicating SPC occurrence. This is consistent with the results of the XRD analysis shown in Fig. 1a, b. The corresponding area fractions of each lateral grain size are shown in Fig. 1i, j. The detected minimum grain size is around 15 nm, which is limited by the electron beam size of the EBSD measurements. For the as-deposited films (Fig. 1i), all films showed the maximum area fraction for a grain size of 15 nm; however, a small proportion of the area fraction with a grain size of~70 nm was detected in the In 2 O 3 film, indicating nuclei in the as-deposited film. After annealing at 250°C (Fig. 1j), the peak of the area fraction shifted toward a larger grain size as R[H 2 ] increased, and the In 2 O 3 :H film deposited at R[H 2 ] = 5% showed a maximum area fraction of 23% at a grain size of 140 nm. Furthermore, as R[H 2 ] increased from 0 to 5%, the area fraction of the minimum grain size below 15 nm significantly decreased, and only a few small grains were in between the large grains, as shown in Fig. 1h. Similar results were observed for films annealed at 250°C in ambient air (shown in Supplementary Fig. 1). The EBSD results show that the nuclei density in the as-deposited film was suppressed by introducing hydrogen during sputtering. Because of the reduction in the nuclei density in the initial In 2 O 3 :H film, the grain size of the In 2 O 3 :H film could be enlarged through SPC. Thus, the XRD and EBSD results indicate that controlling the crystallinity and nuclei density in the as-deposited film are key factors to achieve high-quality In 2 O 3 :H films.   (Fig. 2a), whereas the N e of the In 2 O 3 :H film annealed in air rapidly decreased for T ann > 200°C (Fig. 2c). In addition, the N e reduction was remarkable in the In 2 O 3 :H film deposited at a higher R[H 2 ] value. As a result, an appropriate N e value of 2.0 × 10 17 cm −3 for TFT fabrication was obtained at T ann = 300°C for the In 2 O 3 :H film deposited at R[H 2 ] = 5%; this N e value is over two orders of magnitude lower than that of the In 2 O 3 film deposited without hydrogen introduction (3.0 × 10 19 cm −3 ). Such a large decrease in the N e value of the In 2 O 3 :H films has not been reported before. Although adding H 2 induced the formation of free carriers in the asdeposited films, the N e of the films could be reduced via the relatively low-temperature SPC process and became comparable to that of single-crystalline epitaxial In 2 O 3 films deposited at 650°C (~1 × 10 17 cm −3 ) 36 .

Structural properties of the
Regarding the Hall mobility of the films, In 2 O 3 without hydrogen introduction exhibited an almost constant µ H value over the entire range of investigated T ann values irrespective of the annealing atmosphere, as shown in Fig. 2b R[H 2 ] = 3% and 104.0 cm 2 V −1 s −1 at R[H 2 ] = 5%, indicating that the SPC started at a T ann value between 150 and 200°C. Furthermore, the increased µ H is attributable to the increased grain size, as shown in Fig. 1f-h. As T ann increased, the µ H of In 2 O 3 :H gradually decreased (Fig. 2b). Upon annealing in air (Fig. 2d) To understand the transport properties of the In 2 O 3 :H films after SPC, the relationship between µ H and N e for the films annealed in the range of temperatures between 200°C and 350°C in N 2 and ambient air was summarized, as shown in Fig. 2e. The N e of the In 2 O 3 :H film could be controlled by up to three orders of magnitude. Moreover, for the N e range between 10 19 and 10 20 cm −3 , µ H increases with increasing R[H 2 ], which is attributable to the suppression of grain boundary scattering due to the increasing grain size. For all the films, µ H increased with increasing N e . In general, the grain boundaries have a small impact on µ H in transparent conductive oxides with high N e (>10 20 cm −3 ) because electrons can tunnel through the narrow width (<1 nm) of the grain barriers at high N e values (>10 20 cm −3 ). However, grain boundary scattering is a dominant factor that limits the µ H in films with lower N e 37 . Thus, the observed decrease in µ H with decreasing N e (Fig. 2e) Figure 2f shows the optical absorption (α) spectra of the In 2 O 3 and In 2 O 3 :H films before and after annealing at 250°C in ambient air. The green line represents the In 2 O 3 film deposited without oxygen and hydrogen (only Ar gas), which is shown for comparison. The spectral features that arise as the photon energy exceeds 2.9 eV are due to the absorption associated with the interband transition in In 2 O 3 , whereas the features that arise when the photon energy is below 1.5 eV are due to free carrier absorption. The absorption in the subgap region (<2.9 eV) dropped as R[O 2 ] increased from 0 to 1%, suggesting that oxygen deficiencies, which give rise to subgap defects, are compensated when sputtering in an oxidizing atmosphere. When hydrogen is added during sputter deposition, the absorption is enhanced in the subgap region for the as-deposited films, especially in the photon energy region below 1.5 eV, indicating that free electron absorption is increased due to the hydrogen doping effect. On the other hand, after annealing at 250°C in ambient air (dashed line), the absorption across the subgap of the In 2 O 3 :H film decreased significantly and was lower than that of the In 2 O 3 film. Since subgap defects are generated from native defects, such as oxygen vacancies, as described above, it can be inferred that oxygen vacancies were efficiently reduced in In 2 O 3 :H via SPC in ambient air. This result is consistent with the Hall effect measurements, where it was found that N e decreased from 5.7 × 10 20 to 2.0 × 10 18 cm −3 upon annealing in air at 250°C, as shown in Fig. 2c. The structural, electrical, and optical properties of the In 2 O 3 :H films show that the hydrogen introduced during sputter deposition plays an important role in enlarging the grain size and decreasing the subgap defects after SPC, increasing µ H and decreasing N e . However, a more detailed study will be necessary to carry out quantitative evaluations of the carrier generation and scattering in the SPC-prepared In 2 O 3 :H films.  The variations of µ FE , subthreshold swing (SS), threshold voltage (V th ), and hysteresis (ΔV) were evaluated from ten TFTs on the same substrate. The µ FE was calculated from the linear transfer characteristics as µ FE = Lg m /(WC ox V ds ) at V ds = 0.1 V, where g m is the transconductance, C ox is the oxide capacitance of the gate insulator, and V ds is the drain voltage. V th was defined by gate voltage (V gs ) at drain current (I ds ) of 1 nA, and SS was extracted from V gs , which required an increase in the I ds from 10 to 100 pA. The In 2 O 3 TFT without H 2 introduction (Fig. 3c) Figure 3b shows a cross-sectional conventional bright-field HRTEM image and selective area electron diffraction (SAED) pattern obtained from the SPC-prepared In 2 O 3 :H TFT (R[H 2 ] = 5%). A clear lattice image was observed over the entire thickness of the In 2 O 3 :H channel. Moreover, there was a single crystal-like diffraction pattern in the SAED pattern, even in the thin layers, roughly at a distance of 5 nm from the SiO 2 gate insulator without detectable diffuse ring patterns, which would be attributable to an amorphous phase. This observation explains the high µ FE of 139.2 cm 2 V −1 s −1 in the In 2 O 3 :H TFTs, which is comparable to the µ H of epitaxial single-crystal In 2 O 3 films (~160 cm 2 V −1 s −1 ) 39 . In addition, although In 2 O 3 :H is a polycrystalline film, the standard deviations (σ) of µ FE , SS, V th , and ΔV of the In 2 O 3 :H TFT (R[H 2 ] = 5%) were 3.0 cm 2 V −1 s −1 , 0.02 V dec −1 , 0.2 V, and 0.1 V, respectively, indicating high uniformity of the TFT characteristics (shown in Fig. 3f-i and Supplementary Fig. 2).
To investigate the reliability of the SPC-prepared In 2 O 3 :H TFT (R[H 2 ] = 5%), positive-bias stress (PBS) and negative-bias stress (NBS) tests were carried out under a humidity of 50%. The gate stress voltages for the PBS and NBS were +20 and −20 V, respectively. Figure 4a, Fig. 4b. Furthermore, the V th shift became more significant when the NBS test was conducted at a higher humidity of 70% (shown in Supplementary Fig. 3). Water molecules are coupled to the backchannel of the IGZO TFTs, and excess electrons are donated to the channel under NBS, resulting in a negative V th shift 40,41 . Applying a passivation layer to the TFTs is effective in minimizing the influence of the atmospheric environment; however, hydrogen can diffuse into the channel through the passivation layer and increase the N e of the channel 42 . Although the SiO 2 passive layer was applied to the SPC-grown In 2 O 3 :H TFT, as shown in Fig. 3a, its protection ability was insufficient because the SiO 2 film was deposited via sputtering at RT. Hence, it is believed that the reliability of In 2 O 3 :H TFTs can be improved by selecting the appropriate passivation layer, such as SiN, Al 2 O 3 , or Y 2 O 3 .

Discussion
In this study, we demonstrate the high-performance polycrystalline In 2 O 3 :H TFTs using a low-temperature SPC process.   Characterization of the In 2 O 3 :H films and TFTs. Structural, electrical, and optical measurements were conducted on the 50-nm-thick In 2 O 3 and In 2 O 3 :H films deposited on a synthetic quartz substrate. The films' structural changes were evaluated through X-ray diffraction (XRD) (Philips corp., X'pert) with CuKα radiation and electron backscattering diffraction (EBSD) (EDAX-TSL Hikari High Speed EBSD Detector). The films' carrier concentrations (N e ) and Hall mobility (µ H ) were determined via Hall effect measurements (Accent, HL5500PC) using the van der Pauw geometry at room temperature (RT). The films' optical properties were measured via spectrophotometry (Hitachi, U-4100). The current-voltage characteristics were measured using a semiconductor parameter analyzer (Keysight, E5270B) at RT in the dark. High-resolution transmission electron microscopy (HRTEM) (JEOL, JSM-7001F) analysis was also conducted to observe the microstructure of the In 2 O 3 :H channel in the TFTs.