Improving Electrical Properties by Effective Sulfur Passivation via

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C: Surfaces, Interfaces, Porous Materials, and Catalysis

Improving Electrical Properties by Effective Sulfur Passivation via Modifying Surface State of Substrate in HfO /InP Systems 2

Hang-Kyu Kang, Yu-Seon Kang, Min Baik, Kwang-Sik Jeong, Dae-Kyoung Kim, Jin Dong Song, and Mann-Ho Cho J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b00524 • Publication Date (Web): 20 Mar 2018 Downloaded from http://pubs.acs.org on March 21, 2018

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The Journal of Physical Chemistry

Improving Electrical Properties by Effective Sulfur Passivation via Modifying Surface State of Substrate in HfO2 /InP Systems Hang-Kyu Kang1,2, Yu-Seon Kang1,3, Min Baik1, Kwang-Sik, Jeong1 Dae-Kyoung Kim1 , Jin-Dong Song2, Mann–Ho Cho1,*

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2

3

Department of Physics and Applied Physics, Yonsei University, Seoul, 120-749, Korea

Center of Opto-Electronic Materials, Korea Institute of Science and Technology, Seoul 136-791, Korea

Process Development Team, Semiconductor R&D center, Samsung Electronics Co., Ltd., Yongin 446-711, Korea.

Abstract The thermal stabilities and interfacial properties of HfO2 films created on conditioned i-InP surfaces using atomic layer deposition were investigated.

When HfO2 deposited on sulfur

passivation InP substrate, improved interfacial properties and electrical properties were observed by suppressing the interfacial oxides and In or P dangling bonds between HfO2 and InP.

X-ray

photoelectron spectroscopy (XPS) and thermodynamic data indicated that Acetone-MethanolIsopropanol (AMI) pre-clean process on InP substrate before Sulfur treatment helps the formation of sulfur passivation layer on InP surface more effectively, as comparing to hydrogen fluoride (HF) precleaning.

HF pre-cleaning reduces InP native oxides effectively, but In-F bonds are generated on InP

surface which interrupts the formation of In-S bonds.

Moreover, total density of states (TDOS) and

electron localization functions (ELFs) calculation data showed that In-F bonds does not significantly decrease mid-gap defect states induced by In dangling bond because Fluorine does not chemically bond with In atoms. As a results, AMI pre-clean process were proposed for effective S passivation on substrate in HfO2/InP system. The capacitance–voltage (C–V) data revealed that the hysteresis width and frequency dispersion in the C–V accumulation and depletion were significantly improved in the AMI+S treated sample, as compared with HF+S treated sample. In addition, the AMI+S

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treated HfO2/InP showed excellent thermal stability for the interfacial, structural, and electrical properties during post annealing at 600°C.

* Electronic mail: [email protected]

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I. INTRODUCTION With the world trending towards high-speed applications and lower-power consumptions, SiO2/Si-based metal-oxide-semiconductor (MOS) devices are currently reaching their fundamental limits.1-2 To improve their performance characteristics, many researchers have studied group III–V compound semiconductors that have better carrier transport properties and lower power consumption than that of the Si-based devices.1-3 Because of these advantages, electrical devices using group III–V semiconductors maybe considered the next-generation of Si-based devices.3 Among the group III–V semiconductors, InP has been actively studied due to its high electron mobility and high breakdown endurance.3 Moreover, the downscaling of SiO2 has now reached its physical limits due to critical issues such as poor reliability, high-frequency dispersion, and high current leakage. To overcome these problems, high-κ materials are under serious consideration for replacing SiO2. 2-4 Therefore, recent research on HfO2/InP has been focused on its dielectric properties, thermal stability, and low current leakage.4 However, the thermal diffusion of substrate elements into high-κ device films and poor interfacial qualities between the dielectric and substrate critically affects device instability and leads to device deterioration. To improve interfacial characteristics, many researchers have studied the effect of surface passivation before deposition of the high-κ films.5,6,7 In particular, sulfur passivation using a (NH4)2S treatment has been reported to reduce the interfacial reactions between the high-κ films and the group III–V substrates. In a HfO2/InGaAs system, the sulfur treatment can effectively reduce the interface trap density (Dit) between HfO2 and InGaAs substrates. In this system, the capacitance–voltage (C–V) spectral hysteresis was improved both before and after the annealing process due to the InGaAs surface treatment. 8 Moreover, for GaAs systems, sulfur passivation using (NH4)2S was reported to be effective in suppressing the oxidation states related to Ga and As diffusion, resulting in improvements in the C–V characteristics of frequency dispersions.9,10 Recently, it has been reported that sulfur passivation was successful in reducing the interface defect density, as shown in the decrease in frequency dispersion of C-V in HfO2 /InP before and after post-deposition annealing (PDA) at 600 °C. In this study, in comparison with samples without the sulfur treatment, the Dit and 3



defect states related to out-diffusion of substrate atoms into the HfO2 film were significantly decreased using the sulfur treatment.11 However, many researchers have only considered the process of sulfur treatment, ignoring the effects of the surface cleaning prior to (NH4)2S treatment. A recent study reported that the surface treatment with hydrogen fluoride (HF) cleaning does not maintain the perfect hydrophobic surface state condition of group Ⅲ-Ⅴ substrates.12 Therefore, the formation of In-F bonds can induce the hydrophilicity surface of InP, indicating that HF pre-treatment can hinder the formation of a perfect hydrophobic surface on InP substrates using sulfur passivation by the (NH4)2S solution. Based on these results, we focused on the surface treatment method to enhance sulfur passivation. In this paper, we investigated various surface cleaning methods: i) 1% hydrogen fluoride (aqueous) for 3 min cleaning only, ii) 1% HF for 3 min followed by 21% (NH4)2S cleaning for 10 min, and iii) Acetone-Methanol-Isopropanol (AMI) for 5 min each, sequentially, followed by 21% (NH4)2S cleaning for 10 min. In comparison with the various cleaning methods, the sulfur passivation showed better enhancement of the interfacial stability between the dielectric layer and substrates, which is consistent with other reports. By introducing the AMI surface cleaning followed by 21% (NH4)2S passivation, we observed better surface cleaning effect compared with previous reports on sulfur treatment for effective suppression of interfacial oxides. Significantly improved MOS capacitor characteristics, such as C–V characteristics and Dit, indicate that interfacial reaction before and after the annealing process at 600°C can be effectively controlled by introducing the AMI surface cleaning procedure into the (NH4)2S passivation process.

II. EXPERIMENTAL In general, the HF solution treatment should effectively control the interfacial oxides on group Ⅲ–Ⅴ semiconductors. The combination method of HF treatment followed by (NH4)2S treatment is a well-known method for replacing interfacial oxides with In–S bonds. This typically improves the substrate’s interfacial characteristics and results in stable electrical properties in a high-κ/Ⅲ-Ⅴ 4


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semiconductor system. To optimize the surface passivation, we focused on the AMI pretreatment because this cleaning process does not generate In–F bonds that would hinder S passivation. Before the deposition of the HfO2 films, we applied three different surface cleaning methods: i) Soaking the InP substrates in 1% HF solution (~1% HF in deionized water) for 3 min. ii) Soaking the InP substrates in 1% HF solution for 3 min followed by cleaning with (NH4)2S for 10 min to achieve sulfur passivation. iii) Soaking the InP substrates in acetone, methanol, and isopropanol (AMI) for 5 min each, sequentially, followed cleaning with 21% of (NH4)2S for 10 min. The substrates were then rinsed in deionized (DI) H2O and dried by blowing N2 over the substrates. HfO2 films were deposited on i-InP(001) substrates (doping concentration of ~ 1015 cm−3) using atomic layer deposition (ALD). The chemically etched InP substrates were loaded into the ALD chamber within a few minutes to prevent the formation of native oxides on the substrates. The ALD temperature was 250°C and tetrakis (ethylmethylamino) hafnium (TEMAHf), N2, and H2O vapor were used as the Hf metal precursor, carrier gas, and oxygen source, respectively. The ALD growth process was followed by a N2 purging step: TEMAHf (2 s) → N2 (10 s) → H2O (1.5 s) → N2 (15 s) represents one complete cycle. The flow rates for TEMAHf and H2O were 10 sccm and 50 sccm, respectively, and that of the purging N2 was 200 sccm. The growth rate of H2O in our ALD system was ~ 0.78 Å / cycle. Therefore, we performed the 77 cycles of ALD to deposit the HfO2 film to achieve a film thickness of 6nm. We confirmed a film thickness of 5.7nm using the HR-TEM. After deposition, the samples were annealed at 600°C under N2 at atmospheric pressure for 1 min (PDA process). The microstructure and film thickness of the HfO2 films on the InP substrates were investigated using high-resolution transmission electron microscopy (HR-TEM) with an accelerating voltage of 200 kV. The interfacial chemical bonding configuration was examined using high-resolution X-ray photoelectron spectroscopy (HR-XPS) with a monochromatic Al Kα X-ray source (hv = 1486.7 eV) and a pass energy of 20 eV. The binding energy of the measured core-level spectra of In 3d (Hf 4p), Hf 4f (In 4d), P 2p, C 1s, and O 1s were taken. We calibrated the binding energies of the XPS core5



level spectra using an InP substrate spectrum of 444.4 eV for the InP substrate peak in the In 3d spectra. To deconvolute the chemical states of the XPS core-level spectra, the background was removed by Shirley-type subtraction, and the full width at half-maximum (FWHM) values of the constituent peaks were uniformly fixed. The peaks of the XPS core-level spectra were fitted using a Gaussian–Lorentzian distribution. For the In 3d and P 2p doublets, the intensity ratio of the doublets caused by spin-orbit splitting was determined using the transition probability during the photoionization step. To examine the electrical characteristics of the films, a MOS capacitor (MOSCAP) with a sputter-deposited TiN top electrode (64 × 10−5 cm2 area and 500 nm thickness) was fabricated via the lift-off technique. The C–V curves were measured using an Agilent E4980A LCR meter and an Agilent B1500A semiconductor analyzer. To obtain the Dit of HfO2 /InP at the as-grown and at an annealing temperature of 600°C using rapid thermal annealing (RTA) states, capacitance (Cm) and conductance (Gm) were determined by measuring the parallel conductance (Gp /ω)max. The (Gp /ω)max was calculated using the following equation:12, 13

=

[ ]

Where ω is 2f (f is the measured frequency from 1kHz to 1 MHz). Cox is measured capacitance in accumulation region in C-V measurement, Gm is the measured conductance, and Cm is the measured capacitance. The Dit is subtracted from the peak value of Gp /, using the following relations: 12, 13

D ≈ 2.5

/ "

%$#

,

Where A is the area of the electrode (2.5 × 10-12 cm-2) and q is the elemental charge. 13, 14 To predict the passivation effect of S and F on the InP surface, we performed density functional theory (DFT) calculations with the Vienna ab-initio simulation package (VASP) code. 15, 16 and the 6


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GGA-PBE sol17,

18

exchange correlation function. First, the unit cell of InP was geometrically

optimized with 9 × 9 × 9 k-points and 500 eV cut-off energy until the system energy converged under 0.02 eV/Å. From the optimized unit cell, a (001) surface model consisting of ten In-P layers was constructed. One side of the surface model was terminated with In atoms to reflect a true InP(001) surface, while the other surface was terminated by H atoms with 1.25 charge to simulate the bulk charge configuration of bulk InP. To minimize interactions between each surface, the c-axis of the model was extended to 45 Å using a vacuum slab. Dangling bonds were constructed by having two In atoms approach, and then, by performing geometry optimization with 7 × 5 × 1 k-points and 500 eV cut-off energy until the energy converged below 0.05 eV/Å. The passivated models were also geometry optimized with the same conditions.

III. RESULTS AND DISCUSSIONS The structures of ALD grown HfO2 on the InP substrates prepared with either no treatment (reference), HF cleaning only, HF + (NH4)2S (S) treatment, or AMI + S cleaning before and after PDA at 600°C were investigated using cross-sectional HR-TEM. The thickness of the as-deposited HfO2 film was ~5.7 nm with an amorphous phase. No interfacial layers were observed in all of the as-deposited films as shown in Figure S1 (a)–(d). A clean interface between the HfO2 and InP substrates can be achieved by wet treatment and self-cleaning during the ALD process.19 In general, the TEMAHf precursor effectively reduces the level of native oxides at the interface of high-κ/Ⅲ-Ⅴ semiconductors during few ALD cycles. Moreover, interfacial oxides between the HfO2 and InP are critically affected by the film thickness. That is, the interfacial layer is regrown in thin HfO2 films (< 2 nm) through the diffused oxygen.20 On the other hand, that is nearly perfectly suppressed in relatively thick HfO2 films (> 6 nm).20 Therefore, HfO2 film thickness is not significantly changed even after the PDA process as shown in the HRTEM image of Figure S1. We tried to confirm the change in the interfacial layer using a spectroscopic method such as XPS. In comparison with before and after the PDA at 600℃, the amount of interfacial oxides such as In2O3, In(PO3)3, and InPO4 are 7



not significantly increased, which is consistent with the HR-TEM data In order to investigate the interfacial layer in detail, HR-XPS measurements were performed. Figure 1 shows the In 3d (Hf 4p) and P 2p core-level HR-XPS spectra on the InP substrate as a function of various surface cleaning methods before the deposition of HfO2 films by ALD. The In 3d and P 2p spectra were calibrated with the InP substrate as a reference. On the basis of the phase diagram of an In–P–O ternary system, In2O3 (ΔG ~ −198.6 kcal/mol) at 444.7 eV, In(PO3)3 (ΔG ~ 610 kcal/mol) at 445.4 eV, and InPO4 (ΔG ~ -287 kcal/mol) at 445.7 eV are shown in Figure 1 as the native oxides that may initially be formed on the InP substrates at room temperature. In comparison with the other oxidation states, the In2O3 states were dominantly generated on the InP substrate. At room temperature, these results can be explained by the following possible thermodynamic equilibrium: 4 InP + 3 O2 → 2 In2O3 +4 P (ΔG ~ −323.6 kcal/mol), and 3 InP + 6 O2 → In(PO3)3 + In2O3 (ΔG ~ −753.4 kcal/mol).20 In addition, we observed that the multiple oxidation states can be clearly differentiated between the two S treatments and the HF cleaning only on the InP substrates, as shown in the In 3d and P 2p data. Although the In–S and In–O bonding states cannot be clearly distinguished by HR-XPS because they have nearly identical bonding energies,21 the signal increase after the HF+S and AMI+S treatment can be attributed to the formation of In–S bonds.23 When the InP surface was treated with the (NH4)2S solution, new chemical species may appear as S atoms are formed by bridging bonds to surface In atoms.22 That is, the S elements preferentially react with In elements to form In–S states at the surface of the InP substrate, which suppresses the formation of In–O bonds, as reported by H. S. Jin.23 This is consistent with our In 3d and P 2p spectra; in particular, using the (NH4)2S solution, the change in In(PO3)3 and InPO4 can well support the effective control of oxidation states because of the overwhelming preference of In–S bonds at the surface. Based on the In 3d XPS spectra, the relative ratio of oxidation states to the InP substrate by the various surface treatment methods was derived from the In 3d core-level spectra of Figure 1 (a) and Figure S2. Interestingly, we observed that AMI+S treatment can be more effective in controlling the surface oxidation states in comparison with the HF+S treatment. In particular, the InPO4 states 8


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significantly decreased in the AMI+S treated sample. This was probably due to the influence of different surface cleaning method prior to the S treatment. In general, when an InP substrate is etched using HF solution, In–F bonds form and P elements are changed to PH3 states at the surface of the InP substrates.24 According to the chemical reaction mechanism, when a HF+S treatment is performed, the S atoms decompose the existing In–O bonds to form In–S bonds; however, the presence of In–F states may obstruct the creation of In–S states. In contrast, when AMI surface cleaning was performed before the (NH4)2S treatment, more of the In–O states become converted to the In–S state. Figure 2 (a) and (b) show the In 3d and P 2p core-level HR-XPS spectra for the HfO2 (2 nm) film grown on an InP substrate. Comparing the oxidation states of the reference with those of the HF cleaned only sample, it was found that the oxidation states of In2O3 and In(PO3)3 at the interface of HfO2/InP significantly decreased. This was due to the effective removal of the native oxides on the InP substrate by the HF solution. Furthermore, we confirmed the difference in O 1s spectra depending on the surface treatment methods as shown in S3. Comparing the reference samples with the other samples (HF, HF+S, and AMI+S treated on samples), the intensity of O 1s spectra was significantly reduced in the surface treatment methods. However, in HF+S and AMI+S treatment, the In-P-O bonds are significantly decreased at HfO2 /InP, while the peak caused by In-O state is rather increased compared with HF only treated on samples. These inconsistent results can be explained: i.e., since the In-S and In-O have nearly identical chemical bonding energy21, the peak increase after the sulfur treatment can be reasonably attributed to the formation of In-S bonds.23, 25 In addition, given these results, in HF+S or AMI+S treatments, the intensity of O 1s spectra shows no significant changes compared with HF treatment, as shown in Figure S3. This result is consistent with the peak increase in In 3d shown in the AMI+S-treated sample, which resulted from the increase in the In-S state, not from the In-O state. In a previously reported study, In and P atoms diffuse into the HfO2 film, and the residual oxygen can easily diffuse through the HfO2 film.26 The diffused oxygen can react with the InP substrate to generate other oxidation states. On the contrary, the (NH4)2S solution treatment 9



suppresses the formation of In–P–O bonds between HfO2 and InP because the S passivation process forms In–S states at the interface.27 The formation energy of In–S states compared with that of the In– O states can well explain the suppressed reaction process. For this reason, In–S states effectively control the In–O and In–P–O bonds at the interface, as shown in the HF+S and AMI+S treated samples. Moreover, in the P 2p XPS spectra of the (NH4)2S treated InP substrates, the In–P–O bonds, such as In(PO3)3 and InPO4, were more effectively suppressed at the HfO2/InP interface than that in the reference and HF cleaned only samples. These results are consistent with Figure 2 and Figure S4. To investigate the influence of the various cleaning effects on the electrical properties, the capacitance–voltage (C–V) curves were measured at the frequencies of 10 kHz, 100 kHz, and 1 MHz, as shown in Figure 3 (a)–(d), and with variable frequencies from 1 kHz to 1 MHz, as shown in Figure 3 (e) – (h). The values of the accumulation capacitance are as follows; Cox =1.8 )F/cm2 for as-grown sample, Cox =1.9 )F/cm2 for the HF-treated sample, Cox =2.0 )F/cm2 for HF+S treated sample, and Cox =1.9 )F/cm2 for AMI+S treated sample at room temperature and 100 kHz, respectively. The C–V hysteresis characteristics in the reference and the HF cleaning only samples were very poor. Moreover, in the variable frequency C–V curves, a large frequency dispersion in the accumulation and depletion regions were observed, indicating that the HF cleaning only method did not significantly improve the interface in the HfO2 / InP sample. On the contrary, both the HF+S and the AMI+S treated samples demonstrated significantly decreased hysteresis widths. This results strongly supports the claim that the additional sulfur treatment method directly affects the electrical characteristics. In general, interfacial defects, such as oxygen vacancies and oxygen interstitial sites, influence the charge trapping and de-trapping processes, resulting in a degradation of the C–V characteristics.28 Although the HF surface treatment is an efficient method to remove the oxidized layer of the InP substrate, it is not effective in passivating the defective bonds at the HfO2/InP interface. Therefore, the HF only treated sample has limited control when it comes to the defective interfacial states between the HfO2 and InP, as shown in Figure 3 (b) and (f). In contrast, the additional S treatment can reduce the interfacial state density by passivating the defective bonds with sulfur, resulting in significantly 10


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improved C–V characteristics. Interestingly, the HF+S treated samples and the AMI+S treated on samples showed very different C–V results. In particular, a decrease in the C–V hysteresis width was clearly observed in the AMI+S treated sample, as shown in Figure 3 (c) and (d). In comparison with HF+S treated sample, in the variable frequency C–V data of the AMI+S treated sample, a better frequency dispersion in the accumulation and depletion regions was observed. This can be attributed to the AMI surface cleaning. Additionally, the inversion region of the variable frequency C–V data was also improved in the AMI+S treated sample, as shown in Figure 3 (g) and (h). This result correlates well with the interface trap response rather than with the generation of minority carriers.29 As previously mentioned, during the HF solution etching of InP, In–F bonds are generated, and this releases PH3 as a byproduct at InP substrate surface. The presence of In–F bonds on the InP surface interferes with the formation of In–S on the InP surface. Therefore, the formation of In–F bonds can be a major hindering factor while considering control of the defective interfacial states in HfO2/InP systems. However, the formation of In–S states using the AMI cleaning followed by S passivation method is more complete. Thus, S atoms can more easily diffuse through the oxidized InP surface, resulting in a S passivation that completely controls the defective interfacial states in HfO2 / InP systems.30 Figure 4 shows the MOSCAP C–V characteristics of the HfO2 during various surface treatments after PDA at 600°C: no treatment ((a) and (e)), HF cleaning only ((b) and (f)), HF+S cleaning ((c) and (g)), and AMI+S cleaning ((d) and (h)). The values of the accumulation capacitance are as follows; Cox =1.6 )F/cm2 for as-grown sample, Cox =1.6 )F/cm2 for the HF-treated sample, Cox =1.9 )F/cm2 for HF+S treated sample, and Cox =1.7 )F/cm2 for AMI+S treated sample at PDA 600 ℃ and 100 kHz, respectively. After PDA at 600°C, the accumulation capacitances in all of the samples decrease, as shown in Figure 4. In addition, compared with the as-grown HfO2/InP, we observed that the capacitance value in the accumulation region continuously increases more without saturation. This can be explained from the formation of multiple oxidation states during the PDA process. The oxidation states at the interface can generate several defective interfacial states at the 11



HfO2/InP interface. Comparing the reference sample with the HF treated only sample, no effects (reducing the hysteresis width for C–V accumulation dispersion) are visible, as shown in Figure 4 (b) and (f). Moreover, the C–V accumulation dispersion data in the reference sample and HF cleaning only sample increased significantly, which is consistent with increases in interfacial defect states rather than an increased interfacial layer. When the (NH4)2S treatment was performed, the C-V accumulation dispersion showed no significant changes after the PDA process. On the other hand, after the PDA at 600 ℃, the accumulation dispersions in the C-V curve of the reference sample and the HF-treated sample were significantly higher compared with the reference samples. This is directly related to the improvement of interfacial properties by the formation of In–S states and the reduced number of defective interfacial states. Compared with the HF+S treated sample, even after annealing at 600°C, it is clear that the C–V hysteresis characteristics of the AMI+S treated sample significantly decreased in the depletion region. These results indicate that the combination of AMI surface cleaning and (NH4)2S treatment is very effective in improving the interfacial quality. The border-trap density, with respect to the various surface treatments, was extracted from the difference in the forward and reverse C–V curves at 100 kHz using the following equation: [Crf = │Cr - Cf│], where Cr and Cf are the capacitance densities during the reverse and forward scans, respectively, as shown in Figure 5 (a) as-grown and (b) after PDA at 600°C. The shift in the C-V curve toward a positive voltage during the reverse sweep indicates the substrate-injected electron trapping with a slow response time near the interface of HfO2/InP.31 Because the border traps are affected by the time scale of the measurement, the C-V measurement was performed at a slow sweep rate (~ 0.1 Vs-1) to encompass most of the border traps.32 In general, the border traps with a slow response time are strongly related to the weak depletion and lie close to the flat-band region in C-V measurement. The AMI+S treated sample showed the lowest border-trap density. The border-trap density (before annealing) of the reference sample, HF-cleaned sample, HF+S treated sample, and AMI+S treated sample were ~ 1.3 × 1011 cm−2, ~1.37 × 1011 cm−2, 1.35 × 1011 cm−2, and 0.97 × 1011 cm−2, respectively. Since the depletion region is associated with the interface trap and the 12


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accumulation region is related to the oxide traps in the C–V curve, the multiple oxidation states may significantly affect the trap density, as shown by the XPS and accumulation capacitance. Moreover, the multiple oxidation states cause the flat-band shift to positive voltage.32 Given this results, the maximum peak position of reference and HF treated samples shifted toward a positive voltage at ~ 0.9eV compared with HF + S and AMI+S treated samples. Also, the multiple oxidation states on the film cause the formation of trap densities in the HfO2 layer. According to Y. S Kang, the nitridation effect of HfO2/InP systems is very effective at lowering the border-trap density by direct suppression of oxygen vacancies in the HfO2 film.33 Since the main trap density source in HfO2 is caused by oxygen vacancies in the dielectric film, decreasing more of the unwanted oxidation states using the AMI+S treatment can be effective for decreasing the trap density. The border-trap densities after the annealing process at 600°C show the effect of the AMI+S treatment (Figure 5 (b)). The border-trap density of the reference sample, HF-cleaned sample, HF+S cleaned sample, and AMI+S cleaned sample decreased to ~0.9 × 1011 cm−2, ~1.0 × 1011 cm−2, 0.7 × 1011 cm−2, and 0.6 × 1011 cm−2, respectively. The reduced border-trap density can be explained by the elimination of oxygen vacancies in HfO2 during the post annealing process, which is consistent with the reports that oxygen vacancies can be eliminated during the RTA process in ambient oxygenated conditions.34 Interestingly, it was observed that a peak shift of the border-trap density occurred in the samples before and after the S treatment, as shown in Figure 5 (a) and (b). This can be related to changes in the interface states caused by the different surface treatment methods. In general, the In–O and In–P–O states at the HfO2/InP interface are derived from In or P dangling bonds on the InP surface reacting with oxygen from the interlayer diffusion of air. Treatment with (NH4)2S creates In–S states at the HfO2/InP interface instead of In–O and In–P–O states.35 Moreover, when comparing HF+S cleaning with AMI+S cleaning, we can observe that the change in border-trap density, such as quantity change and peak shift, strongly reflects the surface cleaning effect before the (NH4)2S treatment. Therefore, the improved electrical properties that reduce the border-trap density results from the enhanced InP surface passivation using the AMI+S treatment. 13



We also investigated the Dit of the samples both before and after annealing, as shown in Figure 6 (a) and (b). To obtain Dit under the conduction band edge of InP (ECB edge ~ 1.12eV), the capacitance (Cm) and conductance (Gm) were measured using an i-InP substrate under the following conditions: frequency range from 1 kHz to 1 MHz. The HF cleaning effectively decreases the Dit in the as-grown samples; the maximum Dit of ~ 1.8 × 1013 eV−1 cm2 was obtained in the reference sample, whereas the Dit of ~ 7.5 × 1012 eV−1 cm2 was obtained in the HF-cleaned sample, as shown in Figure 6 (a). As mentioned in our C–V results, the defective interfacial states between HfO2 and InP are stabilized by cleaning the HF surface via the removal of interfacial oxides at the HfO2/InP interface. The lower Dit level in the as-grown HF+S and AMI+S treated samples compared with that of the HF only treated sample results from the control of the interfacial defects using sulfur passivation, which is consistent with the C–V results, HR-XPS results, and Figure S5. Interestingly, the AMI+S treated sample shows a smaller Dit level of ~4.2 × 1012 eV−1 cm2 in comparison with the HF+S treated sample’s Dit level of ~ 5.8 × 1012 eV−1cm2. Thus, the sulfur treatment after AMI cleaning of the InP substrate can also effectively improve the interface characteristics between InP and HfO2. These results are consistent with our contour mapping and 2D interface trap densities (Figure S5). After the PDA at 600°C, the Dit level of the reference sample was much lower than that of the reference sample before PDA, as described in Figure 6 (b). This is consistent with the report that the curing effect was maximized owing to the thermal annealing process between the dielectric film and substrate.36 In the reference sample, the curing effect was maximized owing to the lack of surface treatment between the HfO2 and InP. On the other hand, surface-treated samples such as those with HF, HF+S, and AMI+S treatments do not show significant changes in Dit after the annealing process due to the suppression of interfacial trap sites by the formation of stable chemical bonds such as In-F or In-S states. The AMI+S treated samples show the lowest Dit level of ~4.3 × 1012 eV−1 cm2 after PDA at 600°C. Moreover, 2D interface trap densities and contour mapping results of the AMI+S sample show the lowest values (Figure S6). Thus, AMI+S treatment demonstrates the optimal surface treatment method for effective control of the interfacial and electrical characteristics in HfO2 /InP systems. 14


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Since it is not clear to distinguish the chemical bonding states caused by the formation In-F and In-S using XPS data, To obtain the surface structure of InP after F or S atoms have been incorporated, we performed DFT calculations based on the In or P terminated structure of the i-InP(001) unit surface model, as shown in Table 1. Moreover, to analyze the bonding properties between the In–F and In–S states, the electron localization functions (ELFs) were simulated by DFT. The i-InP (001) surface was simulated using single F and S atom substitutions in In or P dangling bond sites. Table 1 shows the formation energies of the surface states in which a single S and F atom is bonded to the surface of an In or P terminated structure. According to the literature, S or F atoms can change the interfacial properties by bonding to In.37, 38 In particular, the In–S bond is well known to prevent the formation of interfacial oxides between thin films and the InP substrate. Based on these explanations, when only the In-terminated structure was considered, we confirmed that the In–F (−196.4 eV) state was more stable than the In–S (−192.9917 eV) state. This calculation implies that forming In–F bonds is easier than forming In–S bonds; therefore, when the S passivation is performed after the HF treatment, the In–S states cannot displace or replace the already formed In–F states. We performed the total density of states (TDOS) calculations for In dangling bonds (In-), the Fpassivated structure (In-F), and the S-passivated structure (In-S) in the In-terminated InP surface model, as shown in Figure 7 (a). Many gap states were generated within the bandgap of the InP, as indicated by the arrows. Compared with the defect states in the TDOS results of the In-F system, the defect states of the In-S system in the mid-gap region were slightly different. In the In-F system, the gap states causing interfacial defects did not significantly change in comparison with the defect states in In- system; only the peak position of the defect states shifted. This result implies that the In–F bonds do not suppress the defect states generated by dangling bonds; i.e., the bonding of F with In dangling bonds on the InP surface does not completely passivate surface defect states. On the contrary, in the In-S system, the density of gap states was significantly decreased in the mid-gap region, as shown in Figure 7 (a). Therefore, In–S can more effectively suppress the defect states in the mid-gap region compared with In–F states. Figure 7 (b)–(d) show a two-dimensional (2D) cross section with 15



the bonding planes of the In-terminated structure, In–F structure, and In–S structure, respectively. The high ELFs distribution around the surface elements indicates the high density of electrons that do not participate in the chemical bond.39 The high charge density of the surface acts as a source for defect states. Figure 7 (c) shows an incomplete passivation effect through the distribution of electrons that do not participate in chemical bonding. Interestingly, the ELFs distribution suggests that the F atom repels the electrons rather than attracting the surplus electrons, as shown in Figure 7 (c). This result can strongly support our TDOS results for shifting the charge states in the In–F structure. Moreover, the combination of In and F forms anti-bonding states rather than a common covalent bond. As a result, in the In–F system, compared with the In dangling bond state in Figure 7 (b), the electrons not participating in a chemical bond are more distributed around the In elements. Therefore, this result shows that In–F states on the InP surface do not completely passivate the active interface between the dielectric film and substrate. Meanwhile, the In–S state can more effectively passivate the interface because the chemical bonds of In–S effectively attract the electrons that do not participate in chemical bonding, as shown in Figure 7 (d). The In–S chemical bond attracts the surplus electrons around the In atoms, resulting in an effective reduction in the density of electrons not participating in the chemical bonds. This simulation describing the chemical bonding between the S and In dangling bond gives further insight to the well-known S passivation effect.

IV. CONCLUSION In summary, the electrical and interfacial properties of HfO2 films formed on i-InP substrates after various surface cleaning methods (HF washing, HF washing + (NH4)2S treatment, and AMI washing + (NH4)2S treatment) were investigated both before and after a PDA at 600°C. The detailed results confirm that the sulfur passivation achieved using (NH4)2S treatment effectively improved the electrical and interfacial characteristics of the HfO2 films both before and after annealing at 600°C. This result can be explained by the preferential reaction of S with In atoms to form In–S states at the surface of the InP substrate, resulting in the suppression of In–O and In–P–O bonds. Moreover, we 16


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confirmed that the surface cleaning method before the S passivation was an important factor for improving the interfacial and electrical qualities of the HfO2 /InP systems. In comparison with the HF+S treatment, the AMI+S treatment was more effective in controlling the defect states of the interface. Thus, it is possible to effectively control the native oxides, such as In2O3, In(PO3)3, and InPO4, on the InP substrate because sulfur can bind well to the dangling In or P bonds. Therefore, using the AMI+S treatment, we demonstrated an enhancement in interface stability and electrical properties between the HfO2 layer and InP. The importance of this surface treatment technique lies in the improvement of the interfacial properties without using additional processes. In comparison with the previously reported surface cleaning methods for high-κ/Ⅲ–Ⅴ compound semiconductors, the collective results reported herein indicate that the AMI+S treatment method is more effective in improving electrical properties and suppressing interfacial defects. Therefore, we have introduced the sulfur passivation method, which improves the electrical properties of HfO2/InP. In addition, this study presents a new methodology that improves the interfacial characteristics of next-generation InP devices.

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ACKNOWLEDGMENTS This work was partially supported by an Industry-Academy joint research program between Samsung Electronics -Yonsei University. Also, this work was partially supported by the Korea Research Institute of Standard and Science (KRISS) under the Metrology Research Center project.

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ASSOCIATED CONTENT. Supporting Information. The Cross-sectional HR TEM images of 5.7nm-thick no treated on HfO2 /InP samples, 1% of HF treated on HfO2 /InP samples, 1% of HF followed by 21% of (NH4)2S (HF+S) cleaned on HfO2 /InP samples, and Acetone-Methanol-Isopropyl followed by 21% of (NH4)2S (AMI+S) cleaned on HfO2 /InP samples for as-deposited and after PDA at 600 ℃; Ratio of oxidation states/substrate by the various surface treatment was calculated based on In 3d core-level spectra in Figure 1 (a); HR-XPS In 3d core-level spectra with takeoff angle of 20º and 90º for reference (no treated on HfO2/InP), HF treated on HfO2/InP, HF+S treated on HfO2/InP, AMI+S treated on HfO2/InP before and after the PDA at 600℃; Normalized parallel conductance (Gp /qA) in MOS capacitor of reference HfO2/InP, HF treated HfO2/InP, HF+S treated HfO2/InP, and AMI+S treated HfO2/InP at room temperature and after the PDA at 600℃. Two-dimensional (2D) interface trap density of reference HfO2/InP, HF treated HfO2/InP, HF+S treated HfO2/InP, and AMI+S treated HfO2/InP before and after the PDA at 600℃.

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