Influence of Grain Boundary Scattering on the Field-Effect Mobility of Solid-Phase Crystallized Hydrogenated Polycrystalline In₂O₃ (In₂O₃:H)

Abstract

Hydrogenated polycrystalline In₂O₃ (In₂O₃:H) thin-film transistors (TFTs) fabricated via the low-temperature solid-phase crystallization (SPC) process with a field-effect mobility (μ_FE) exceeding 100 cm² V⁻¹ s⁻¹ are promising candidates for future electronics applications. In this study, we investigated the effects of the SPC temperature of Ar + O₂ + H₂-sputtered In₂O₃:H films on the electron transport properties of In₂O₃:H TFTs. The In₂O₃:H TFT with an SPC temperature of 300 °C exhibited the best performance, having the largest µ_FE of 139.2 cm² V⁻¹ s⁻¹. In contrast, the µ_FE was slightly degraded with increasing SPC temperature (400 °C and higher). Extended X-ray absorption fine structure analysis revealed that the medium-range ordering in the In₂O₃:H network was further improved by annealing up to 600 °C, while a large amount of H₂O was desorbed from the In₂O₃:H films at SPC temperatures above 400 °C, resulting in the creation of defects at grain boundaries. The threshold temperature of H₂O desorption corresponded well with the carrier transport properties; the µ_FE of the TFTs started to deteriorate at SPC temperatures of 400 °C and higher. Thus, it was suggested that the hydrogen remaining in the film after SPC plays an important role in the passivation of electron traps, especially for grain boundaries, resulting in an enhancement of the µ_FE of In₂O₃:H TFTs.

1. Introduction

Amorphous oxide semiconductors (AOSs), represented by amorphous In–Ga–Zn–O (a-IGZO), have now become standard channel materials in thin-film transistors (TFTs) for active-matrix liquid-crystal displays and active-matrix organic light-emitting diode displays [1,2,3]. This is because a-IGZO has properties that are superior to hydrogenated amorphous Si (a-Si:H), such as a large field-effect mobility (μ_FE) of over 10 cm² V⁻¹ s⁻¹, an extremely low leakage current, a low processing temperature (<350 °C), and excellent uniformity [4,5,6,7,8,9,10,11,12]. Although the μ_FE value of a-IGZO TFTs is more than ten times higher than that of a-Si:H TFTs (<1 cm² V⁻¹ s⁻¹), the further improvement of μ_FE values is required to expand their range of applications as alternatives to poly-Si TFTs (50–100 cm² V⁻¹ s⁻¹) [13]. The optimization of AOS composition is one approach for improving μ_FE; for example, an increase in the In ratio due the considerable spatial spread of the In 5s orbital with a large overlap can provide a facile electron transport path with a low electron effective mass [14,15]. Although various compositions, including In–Sn–Zn–O [16,17], In–W–Zn–O [18,19], Al–In–Sn–Zn–O [20], and In–Ga–Zn–Sn–O [21], have been proposed to enhance the μ_FE, the value remains insufficiently high to compete with that of low-temperature polysilicon (LTPS) TFTs [22].

In contrast, the crystallization of OSs is an another approach for improving μ_FE, because the subgap density of states originating from structural disorder and defects can be suppressed via lattice ordering. Although polycrystalline Oss, including In₂O₃, ZnO, and SnO₂, have been investigated as channel materials in early oxide-based TFTs [23,24,25], they can easily create oxygen vacancies, leading to degenerate semiconductors. In addition, µ_FE degradation due to grain boundary scattering is a serious issue for polycrystalline OSs as well as poly-Si [26,27]. Polycrystalline In₂O₃ films have been investigated for use as the transparent conductive oxide (TCO) in solar cells. Koida et al. reported a degenerate hydrogen-doped polycrystalline In₂O₃ (In₂O₃:H) film with high electron mobility (100–130 cm² V⁻¹ s⁻¹) produced by solid-phase crystallization (SPC) [28]. Recently, we reported a μ_FE value of 139.2 cm² V⁻¹ s⁻¹ for a TFT obtained using hydrogenated polycrystalline In₂O₃ (In₂O₃:H) formed via SPC at 300 °C [29]. The obtained μ_FE value is comparable to the Hall mobility of single-crystalline epitaxial In₂O₃ films (~160 cm² V⁻¹ s⁻¹) [30]. The as-deposited amorphous In₂O₃:H was converted into a polycrystalline film with lateral grain sizes of about 140 nm via SPC [29]. However, the effects of SPC temperature on the electrical and structural properties of In₂O₃:H films are not yet understood in detail.

In this study, we investigated the effects of SPC temperature on electron transport properties in In₂O₃:H films and TFTs. Hydrogen intentionally doped during sputtering was found to play an essential role in the passivation of defects, especially for the grain boundaries of the films, resulting in an enhanced μ_FE of the TFTs.

2. Materials and Methods

2.1. Fabrication of In₂O₃:H TFTs

In₂O₃:H TFTs were fabricated on a heavily doped p-type Si substrate with a 100 nm thick thermally grown SiO₂ layer. The doped p-type Si substrate and SiO₂ layer were used as the gate electrode and gate insulator, respectively. The 30 nm thick In₂O₃:H channels were deposited via pulsed direct current (DC) magnetron sputtering, without substrate heating, from a ceramic In₂O₃ target using a mixture of Ar, O₂, and H₂ gases. The gas flow ratios of O₂ (R[O₂] = O₂/(Ar + O₂ + H₂)) and H₂ (R[H₂] = H₂/(Ar + O₂ + H₂)) were 1% and 5%, respectively. The deposition pressure and DC power were maintained at 0.6 Pa and 50 W, respectively. The base pressure before gas introduction was below 6 × 10⁻⁵ Pa. After deposition, the In₂O₃:H films were annealed in ambient air at a temperature range of 200–600 °C for 1 h. After annealing, a 100 nm thick SiO₂ film was deposited via reactive sputtering without substrate heating. Subsequently, Al source/drain electrodes were deposited via sputtering. Finally, the In₂O₃:H TFTs were annealed at 250 °C in ambient air for 1 h. In₂O₃:H, SiO₂, and Al films were deposited through a shadow mask, and both the channel length and width were 300 µm.

2.2. Characterization Methods

The carrier concentration (N_e) and Hall mobility (µ_FE) of the In₂O₃:H films were determined by Hall effect measurements (Accent, HL5500PC) using the van der Pauw geometry at room temperature. The local structural changes of the films were evaluated through extended X-ray absorption fine structure (EXAFS) at the BL01B1 beamline in SPring-8. The In K-edge fluorescence XAFS of the films was measured using a 19-element Ge detector with an incident X-ray angle of 2° with respect to the sample surface. The XAFS data were analyzed using the Demeter software packages [31]. The macroscopic structure of the In₂O₃:H films was observed using electron backscattering diffraction (EBSD) (EDAX-TSL Hikari High-Speed EBSD Detector). The hydrogen concentration in the films was measured using secondary ion mass spectrometry (SIMS) (ULVAC-PHI, ADEPT-1010) with Cs⁺ as a primary ion. The chemical bonding states of the constituent elements and hydrogen were evaluated from their thermal effusion using thermal desorption spectroscopy (TDS), which was carried out while varying the stage temperature from 50 to 800 °C at a heating rate of 60 °C min⁻¹. Reference films of In₂O₃ under similar conditions maintaining an R[H₂] of 0% (without hydrogen introduction) were deposited for comparison.

3. Results and Discussion

3.1. Electrical Properties of In₂O₃:H Films

Figure 1a shows the variations in N_e and µ_H in the 50 nm thick In₂O₃:H films as a function of the annealing temperature (T_ann). The N_e of the as-deposited film (5.7 × 10²⁰ cm⁻³) began to decrease rapidly from T_ann = 200 °C, where SPC occurred. Then, the In₂O₃:H exhibited an almost constant N_e of ~2 × 10¹⁷ cm⁻³ (an appropriate value for TFT fabrication) over a T_ann range of 300–500 °C. After reaching a T_ann of 600 °C, the N_e slightly increased to 1.4 × 10¹⁸ cm⁻³. The µ_H of In₂O₃:H increased to 78.6 cm² V⁻¹ s⁻¹ after SPC occurred (T_ann = 200 °C), whereas the µ_H of the films decreased to ~15 cm² V⁻¹ s⁻¹ at a T_ann range of 250–350 °C. As T_ann further increased, the µ_H of In₂O₃:H began to decrease, resulting in a µ_H of 0.4 cm² V⁻¹ s⁻¹ at T_ann = 600 °C.

To understand the carrier transport properties of the In₂O₃:H films, the relationship between µ_H and N_e for the films with various T_ann values is shown in Figure 1b. The changes in electrical properties could be classified into the following three regions: (I) enhanced µ_H at a T_ann ≤ 200 °C; (II) decreased µ_H with decreasing N_e at T_ann = 250–350 °C; and (III) decreased µ_H with constant or increasing N_e at T_ann ≥ 400 °C. The decrease in the µ_H of the polycrystalline In₂O₃:H films with increasing T_ann in regions II and III was considered to be due to the effects of grain boundary scattering and intragrain scattering. In general, for degenerate transparent conducting oxide materials, the mobility in the grains is determined by an optical method using the Drude model, because optical mobilities are not affected by grain boundary scattering [32,33]. However, it is difficult to determine the optical mobility of In₂O₃:H films annealed at T_ann ≥ 250 °C because the free electrons are significantly decreased to the order of 10¹⁷ cm⁻³ and the films become non-degenerate semiconductors [34,35]. Therefore, we evaluated the effects of intragrain scattering by measuring the field-effect mobility of In₂O₃:H TFTs. This is because when a voltage is applied to the gate, a large number of carriers (10¹⁹–10²⁰ cm⁻³) are accumulated at the In₂O₃:H/gate insulator interface, and the effects of grain boundary scattering can almost be neglected since electrons tunnel through the narrow width (<1 nm) of the grain barriers at high N_e values [36].

3.2. Electrical Properties of In₂O₃:H TFTs

Figure 2 shows the typical characteristics of the In₂O₃:H TFTs annealed at T_ann values of 200–600 °C. The µ_FE was calculated from the linear transfer characteristics (V_ds = 0.1 V) using the following equation:

$${\mu }_{\mathrm{FE}}=\frac{L{g}_{\mathrm{m}}}{W{C}_{\mathrm{ox}}{V}_{\mathrm{ds}}}$$

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 a 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 average values and standard deviations (σ) of the characteristics of five TFTs on the same substrate are shown in Supplementary Figure S1 and Supplementary Table S1. The In₂O₃:H TFT annealed at 200 °C did not exhibit any switching (conductive behavior) because the In₂O₃:H film was still in a degenerated state (N_e = 5.7 × 10²⁰ cm⁻³). The TFTs annealed at 300 °C exhibited the best performance, with the largest field-effect mobility (µ_FE) of 139.2 cm² V⁻¹ s⁻¹ and smallest subthreshold swing (SS) of 0.19 Vdec⁻¹ with an appropriate threshold voltage (V_th) of 0.2 V (shown in Figure 2c,e). Although the µ_H of the In₂O₃:H films decreased to 14.9 cm² V⁻¹ s⁻¹ with a N_e of 2.0 × 10¹⁷ cm⁻³ after annealing at 300 °C, as shown in Figure 1, extremely high µ_FE values were obtained from the TFTs at higher gate voltages (Figure 2b). Thus, we concluded that the decrease in µ_H in region II shown in Figure 1b) was mainly due to an increase in the potential barrier at the grain boundary caused by a decrease in N_e, rather than a decrease in intragrain mobility. In contrast, the TFT characteristics were slightly degraded when T_ann was increased to 600 °C, i.e., the µ_FE decreased and the SS value increased. Comparing µ_H and µ_FE after T_ann = 600 °C, the µ_H significantly deteriorated to 0.4 cm² V⁻¹ s⁻¹, whereas the µ_FE of the TFTs was maintained at 94.6 cm² V⁻¹ s⁻¹. This result suggests that grain boundary scattering is a dominant factor that limits the µ_H in films annealed at 400 °C and higher.

3.3. Structural Properties of In₂O₃:H Films

To investigate the crystallinity of the films, the effect of T_ann on the local structure of the In₂O₃:H films was evaluated using XAFS. Figure 3 shows Fourier-transformed (FT) EXAFS spectra of the In K-edge for the (a) In₂O₃ and (b) In₂O₃:H films as a function of the phase uncorrected interatomic distance. The as-deposited In₂O₃ film without hydrogen introduction during sputtering exhibited three obvious peaks (Figure 3a), which corresponded to the nearest oxygen (In–O) and the second and third nearest In (In–In and In–In^*). Using the values for the crystalline In₂O₃ powder standard, the interatomic distance (R) and Debye–Waller factor (σ²) for the films were determined. The k-range of the EXAFS data used in the analyses was k = 3–14 Å⁻¹ with a k-weight of 3. The fitting carried out in the R space was from R = 1.0–4.0 Å for the three-shell model. As shown in

Table 1. EXAFS fitting results for first, second, and third shells in the In₂O₃ and In₂O₃:H films.

		1st Shell (In–O)		2nd Shell (In–In)		3rd Shell (In–In*)
	Tann (°C)	RIn–O (Å)	σ2In–O (Å2)	RIn–In (Å)	σ2In–In (Å2)	RIn–In* (Å)	σ2In–In* (Å2)
In2O3	as-depo.	2.16	0.0090	3.35	0.0065	3.82	0.0066
In2O3	300	2.16	0.0073	3.36	0.0062	3.83	0.0063
In2O3:H	as-depo.	2.13	0.0112	-	-	-	-
In2O3:H	200	2.16	0.0072	3.36	0.0059	3.83	0.0051
In2O3:H	300	2.17	0.0072	3.36	0.0056	3.84	0.0049
In2O3:H	600	2.17	0.0072	3.37	0.0051	3.84	0.0046

, the R_In–O, R_In–In, and R_In–In* values of the as-deposited In₂O₃ film without hydrogen were 2.16, 3.35, and 3.82 Å, respectively, which agreed well with the interatomic distance of the In₂O₃ bixbyite structure (space group Ia3, number 206) [37,38,39]. When the In₂O₃ film was annealed at 300 °C, no noticeable changes in R were observed (Figure 3a), while the peak intensity increased in each spectrum, resulting in a decrease in σ². This result indicates that thermal annealing improved the structural disorder of the films. By introducing hydrogen during sputtering, the second and third nearest peaks disappeared in the as-deposited In₂O₃:H film, while the intensity of the first nearest peak decreased (Figure 3b), resulting in an increase in σ²_In–O to 0.0112. In contrast, after annealing at 200 °C, the intensities of all peaks increased significantly and the intensities of the second and third nearest peaks were higher than those of the In₂O₃ film annealed at 300 °C. These results indicate that medium-range ordering was lost around In in the initial In₂O₃:H film, whereas medium-range ordering at distances equal to or longer than the second neighbor significantly improved after annealing at 200 °C. This is in agreement with a previous study using electron backscatter diffraction which confirmed that the amorphous state of the initial In₂O₃:H film and the grain size of the In₂O₃:H film were enlarged through SPC [29]. Therefore, the increase in µ_H after annealing at 200 °C, as shown in Figure 1b) (region I), is due to an increase in grain size as well as the improvement of the local structural order of the In₂O₃:H films. As T_ann increased from 200 to 300 °C, the intensity of the first nearest peak was constant, while the intensities of the second and third nearest peaks increased slightly, resulting in a decrease in σ²_In–In (0.0056) and σ²_In–In (0.0049). This result indicates that the medium-range ordering was improved by annealing at 300 °C, which is in good agreement with the high intragrain mobility obtained via µ_FE for the In₂O₃:H TFT annealed at 300 °C. As T_ann increased from 300 to 600 °C, the intensities of the second and third nearest peaks further increased, whereas R remained almost constant. Despite the improvement in the crystallinity of the In₂O₃:H films observed when annealing at 600 °C, the µ_FE of the TFTs slightly decreased to 94.6 cm² V⁻¹ s⁻¹, as shown in Figure 2c, and the µ_H significantly deteriorated to 0.4 cm² V⁻¹ s⁻¹. The deterioration of the µ_H of the In₂O₃:H films (Figure 1b, Region III) and the µ_FE of the TFTs annealed at ≥400 °C could not be explained by local structural changes in the In₂O₃:H films. Thus, it was suggested that the decreases in the µ_H and the µ_FE of the TFTs in region III shown in Figure 1b and Figure 2c were mainly due to the formation of defects at grain boundaries.

To investigate the origin of the deterioration of the µ_H of In₂O₃:H films at T_ann ≥ 400 °C, we performed EBSD measurements. Figure 4a–e depicts the EBSD images along the normal direction for the In₂O₃:H films with various T_ann values. The as-deposited amorphous In₂O₃:H film was converted into a polycrystalline In₂O₃:H film with grain structure embedded in the amorphous matrix at a T_ann of 150 °C (Figure 4b). At a T_ann of 200 °C (Figure 4c), the film was fully crystallized with a grain size of around 200 nm. After that, no significant difference was observed in the crystal grain size with increasing T_ann values up to 600 °C. The corresponding area fractions of the crystalline phase are shown in Figure 4f. All films showed a maximum area fraction for a grain size of ~200 nm; however, a small proportion of the area fraction with a grain size of ~15 nm was increased in the In₂O₃:H film annealed at 600 °C, as shown in the insets of Figure 4f. Moreover, these small domains were located in between the large grains, as shown in Figure 4e. The results indicate that when the In₂O₃:H films were annealed at 400 °C and higher, small domains were created at grain boundaries that served as electron traps, resulting in a decrease in the µ_H in region III shown in Figure 1b.

To understand the mechanism of structural deterioration at the grain boundaries of In₂O₃:H films at T_ann ≥ 400 °C, we performed TDS measurements. Figure 5 shows the TDS spectra of H₂, H₂O, O₂, and In for the In₂O₃ and In₂O₃:H films. We first note that H₂, O₂, and In desorption were negligible for both types of films, while the H₂O desorption was high for the In₂O₃:H film in particular. Large amounts of H₂O were desorbed at a stage temperature of 400–800 °C for the In₂O₃:H film (Figure 5b). The amount of hydrogen in the film was quantitatively evaluated using SIMS, and it was found that 2.6 × 10²¹ cm⁻³ of hydrogen remained in the film after SPC at 300 °C, which was one order of magnitude higher than that of the In₂O₃ film. The H₂O desorption temperature (400 °C) for the In₂O₃:H film corresponded well to a T_ann of 400 °C, at which the µ_H of the films started to degrade. During annealing at a T_ann of 400 °C and higher, H or –OH inside grains may migrate to grain boundaries and react with a neighboring H at the boundary, resulting in the generation of H₂O molecules. As a consequence, defects are formed at grain boundaries. In other words, the presence of H and/or –OH bonds in the In₂O₃:H film after SPC plays an important role in the passivation of defects, especially for grain boundaries. In general, the SS value of a TFT is strongly affected by defects near the semiconductor Fermi level [2,12]. We recently reported from hard X-ray photoelectron spectroscopy analysis that intentionally introduced hydrogen is effective in reducing defects near the Fermi level in amorphous IGZO [40,41,42,43,44]. The SS values of the In₂O₃:H TFTs, shown in Figure 2d, increased at a T_ann of 400 °C and higher, indicating defect creation. Thus, we believe that the H and/or –OH bonds remained in the films after SPC passivated the defects near the Fermi level, leading to the high µ_FE and steep SS values of the In₂O₃:H TFT annealed at 300 °C. Thus, it is worth noting that grain boundary scattering, which is a serious issue for polycrystalline Si TFTs, may not have a strong influence on the µ_FE of polycrystalline In₂O₃:H TFTs.

4. Conclusions

In summary, we investigated the effects of annealing temperature on the electron transport properties of In₂O₃:H films and TFTs. The changes in the electrical properties of the In₂O₃:H films were classified into the three following regions.

(I) When T_ann = 200 °C, µ_H increased by converting amorphous In₂O₃:H into a polycrystalline In₂O₃:H film with an increase in grain size.

(II) When T_ann = 250–350 °C, µ_H decreased with decreasing N_e. However, when µ_H exhibited low values, the medium-range ordering in the grains improved. The In₂O₃:H TFT annealed at 300 °C exhibited the best performance, with a µ_FE of 139.2 cm² V⁻¹ s⁻¹.

(III) When T_ann ≥ 400 °C, although µ_H significantly decreased with constant or increasing N_e, the µ_FE of the TFTs was maintained at 94.6 cm² V⁻¹ s⁻¹. The medium-range ordering of the In₂O₃:H network was improved by a T_ann of 600 °C, while a large amount of hydrogen was desorbed at 400–800 °C, resulting in defect creation at the grain boundaries. Thus, it was suggested that the hydrogen remaining in the film after SPC plays an important role in the passivation of electron traps, especially for the grain boundaries of the In₂O₃:H films, resulting in an enhancement of the µ_FE. We believe that the SPC-grown In₂O₃:H TFTs are promising candidates for use in future electronics applications.