Reset First Resistive Switching in Ni1−xO Thin Films as Charge Transfer Insulator Deposited by Reactive RF Magnetron Sputtering

Reset-first resistive random access memory (RRAM) devices were demonstrated for off-stoichiometric Ni1−xO thin films deposited using reactive sputtering with a high oxygen partial pressure. The Ni1−xO based RRAM devices exhibited both unipolar and bipolar resistive switching characteristics without an electroforming step. Auger electron spectroscopy showed nickel deficiency in the Ni1−xO films, and X-ray photoemission spectroscopy showed that the Ni3+ valence state in the Ni1−xO films increased with increasing oxygen partial pressure. Conductive atomic force microscopy showed that the conductivity of the Ni1−xO films increased with increasing oxygen partial pressure during deposition, possibly contributing to the reset-first switching of the Ni1−xO films.


Introduction
Resistive random access memory (RRAM) [1] has been widely studied as a candidate for next-generation non-volatile memory to overcome the limitations of conventional memories, such as flash memory and dynamic random access memory (DRAM). RRAM has a relatively low operation voltage with excellent program and erase speed [2]. In addition, the device could be fabricated in a simple metal-insulator-metal (MIM) [3] structure, enabling the high-density cell structure of a cross-bar array with 4F 2 [4,5]. It was reported that numerous transition metal oxides, including Al 2 O 3 [6,7], HfO 2 [8][9][10], NiO x [11][12][13][14], TiO x [15,16], TaO x [17,18], Nb 2 O 5 [19,20], and Pr 1−x Ca x MnO 3 [21][22][23] show resistive switching (RS) characteristics. Moreover, various deposition techniques, such as sputtering [24][25][26][27][28], atomic layer deposition (ALD) [29] and pulsed laser deposition (PLD) [30] were used for the formation of such oxides. Notably, nickel oxide (NiO) film is one of the most widely studied oxides and is reported to have low operation power, a high on/off resistance ratio and is compatible with the CMOS fabrication process [31,32]. NiO has a rock salt structure composed of Ni 2+ and O 2− and is a member of the strongly correlated 3d transition metal oxides that exhibit charge-transfer insulator behavior [33,34]. It is an insulating oxide with a wide bandgap (E g ≈ 4.3 eV) due to the charge transfer gap caused by "Hubbard U" between the 2p and 3d states [34,35]. Therefore, the pristine state of NiO is typically the insulating state in RRAM [36,37]. The RS phenomenon in NiO has been mainly described as the formation and rupture of conductive filaments. This reversible resistance transition between the high-resistance state (HRS) and low-resistance state (LRS) is caused by applying electrical stress after an "electroforming" step [38]. It was suggested that oxygen atoms are migrated by the electric field, leaving oxygen vacancies (V o 2+ ) at the vacated sites during the electroforming step; the adjacent Ni 2+ atoms are changed to Ni 0 to compensate for the charge state, resulting in a Ni filament [39][40][41]. The electroforming process degrades the chemical and physical properties of devices of MIM structure, affecting their reliability. The characteristics of RS uniformity also deteriorate because of non-uniform filament formation among MIM devices [42]. Moreover, electroforming requires additional high-voltage circuits, significantly reducing the device density. Therefore, research on devices that can be operated without an electroforming step is essential for realizing RS memories [43][44][45].
This study investigated the RS characteristics of off-stoichiometric Ni 1−x O films for unipolar and bipolar RSs (URS and BRS, respectively). Particularly, it was demonstrated that nickel-deficient Ni 1−x O films deposited under excessive oxygen partial pressure exhibit a reset-first RS without an electroforming step. An RRAM device with a reset-first RS could be an alternative to overcome the limitations of RRAM requiring an electro-forming step.

Experimental
MIM devices with Pt/NiO/Pt and Pt/NiO/TiN stacks were fabricated for electrical characterization. First, Ti/TiN adhesive layers with thicknesses of 10-50 nm were deposited onto SiO 2 on a Si substrate using DC magnetron sputtering. Pt or TiN films were then deposited as bottom electrodes (BE). BE with various areas of 0.18~4.0 µm 2 were formed to investigate the area-dependence of the electrical characteristics. After BE formation, off-stoichiometric Ni 1−x O films with a thickness of 10 nm were deposited via reactive RF magnetron sputtering using a Ni target under various O 2 partial pressures. During sputtering, the base and working pressures were less than 3 × 10 −3 and 3 mTorr, respectively. During deposition, the RF power and temperature of the substrate were main-tained at 100 W and 400 • C, respectively. The fraction of the O 2 partial pressure in the mixture of Ar and O 2 varied from 10% to 50% for deposition. Finally, Pt top electrodes (TEs) with a thickness of 100 nm were formed using DC magnetron sputtering and a lift-off process. The electrical characteristics of the device were characterized using a Keysight B1500A analyzer at 21~23 • C. RS under DC bias was measured with a com-pliance current of 10 mA to avoid hard breakdown of the Ni 1-x O films. The spatial distribution of conductivity in the pristine state was investigated using conductive atomic force microscopy (C-AFM) (Park Systems, XE-100) with a measurement bias of 3 V [46,47]. Grazing incidence X-ray diffraction (GI-XRD, Rigaku SmartLab), Auger electron spectroscopy (AES, PHI-700, ULVAC-PHI), and X-ray photoelectron spectroscopy (XPS, K-alpha, Thermo U. K.) analyses were conducted to investigate the crystallinity, composition, and valence states of Ni in the Ni 1−x O films, respectively.

Results and Discussion
XRD analysis was conducted to investigate the crystallinity of Ni 1−x O films. The XRD patterns of Ni 1−x O films deposited under various O 2 fractions are illustrated in Figure 1a. The peaks of NiO (111), NiO (200), NiO (220), and NiO (311) imply a polycrystalline structure [48]. NiO films, deposited with an O 2 partial pressure fraction of 50% showed lower intensity with a more comprehensive full-width half maximum (FWHM), implying poorer crystallinity of NiO films. The XRD peak of the (111) plane shifted to lower diffraction with increasing O 2 partial pressure, indicating an increase in the lattice constant with increasing O 2 partial pressure, as shown in Figure 1b. The increase in the lattice constant could be ascribed to the increased strain effect as Ni vacancies increase with excessive O 2 partial pressure [48][49][50]. Figure 1c shows Figure 2a shows the typical behavior of Pt/Ni1−xO/Pt stacks. The pristine Ni1−xO film deposited under an O2 partial pressure fraction of 10% offered an initial high resistivit [51] at an applied voltage of 1.77 V (1.4 MV/cm) on the TE. The film resistance change from HRS to LRS during the forming step. The resistance state was changed back to HR at 0.64 V (0.5 MV/cm) during the subsequent bias application, exhibiting reversible switch ing for the positive bias on TE. The difference between the forming voltage (Vform) and se voltage (Vset) was approximately 0.57 V. In contrast, pristine Ni1−xO films deposited unde the 30% or 50% O2 ratio showed low resistance in the pristine state without the electro forming step and reset-first RS behavior, where the initial LRS state was changed to th HRS state, as shown in Figure 2b,c. While Vset is similar to that of Ni1−xO films for the O partial pressure fraction of 10%, the IHRS/ILRS ratio decreased because of the overall hig current level in the HRS state. In particular, the IHRS between these oxygen partial pressur fractions showed that the 50% O2 ratio was 10 times higher than that of 30% O2. The Icurves of TiN/Ni1−xO/Pt stacks are plotted in Figure 2d-f. The Ni1−xO film deposited unde a 10% O2 partial pressure fraction show BRS [52] characteristics, as shown in Figure 2d The pristine Ni1−xO film showed high resistivity, and the resistance state changed to LR after the electroforming step with a negative bias on TE. The difference between Vform (−4 V) and Vset (−0.7 V) was approximately 3.3 V. On the contrary, the Ni1−xO film deposite under the 30% or 50% O2 partial pressure fraction showed reset-first BRS behavior for positive voltage on the TE, as shown in Figure 2e,f.   Figure 2a shows the typical behavior of Pt/Ni 1−x O/Pt stacks. The pristine Ni 1−x O films deposited under an O 2 partial pressure fraction of 10% offered an initial high resistivity [51] at an applied voltage of 1.77 V (1.4 MV/cm) on the TE. The film resistance changed from HRS to LRS during the forming step. The resistance state was changed back to HRS at 0.64 V (0.5 MV/cm) during the subsequent bias application, exhibiting reversible switching for the positive bias on TE. The difference between the forming voltage (V form ) and set voltage (V set ) was approximately 0.57 V. In contrast, pristine Ni 1−x O films deposited under the 30% or 50% O 2 ratio showed low resistance in the pristine state without the electroforming step and reset-first RS behavior, where the initial LRS state was changed to the HRS state, as shown in Figure 2b,c. While V set is similar to that of Ni 1−x O films for the O 2 partial pressure fraction of 10%, the I HRS /I LRS ratio decreased because of the overall high current level in the HRS state. In particular, the I HRS between these oxygen partial pressure fractions showed that the 50% O 2 ratio was 10 times higher than that of 30%  Figure 3b shows the electrical currents at 0.64 V in the LRS states, which has a similar tendency to the I HRS with O 2 partial pressure, but the slope was lower than that of the I HRS state. The I HRS and I LRS showed the highest values for Ni 1−x O films deposited under the 50% O 2 partial pressure fraction.
To understand the nature of resistance switching, HRS and LRS resistances were measured from devices with BE of 0.18, 0.38, 2.00, and 3.69 µm 2 at a bias of ±0.48 V. Figure 4a shows the area dependent resistance for BRS device with Ni 1−x O films deposited by 10% O 2 partial pressure fraction. The resistance of the HRS remained almost constant with decreasing geometric device area, while that of the LRS is almost independent of the device area. These area-independent characteristics imply that resistance switching through the device occurs in local regions, such as filament paths, rather than homogeneously distributed switching paths [53][54][55][56][57]. Meanwhile, the resistances of reset-first RS devices with Ni 1−x O films deposited at 50% O 2 partial pressure showed increased dependence on the device area, as shown in Figure 4b. Because the area dependence of the LRS for Ni 1−x O films with 50% O 2 partial pressure is close to that of Ni 1−x O films with 10% O 2 partial pressure, the nature of the RS is filamentary in the local area. The significant dependence fractions showed that the 50% O2 ratio was 10 times higher than that of 30% O2. Th curves of TiN/Ni1−xO/Pt stacks are plotted in Figure 2d-f. The Ni1−xO film deposited u a 10% O2 partial pressure fraction show BRS [52] characteristics, as shown in Figur The pristine Ni1−xO film showed high resistivity, and the resistance state changed to after the electroforming step with a negative bias on TE. The difference between Vform V) and Vset (−0.7 V) was approximately 3.3 V. On the contrary, the Ni1−xO film depo under the 30% or 50% O2 partial pressure fraction showed reset-first BRS behavior positive voltage on the TE, as shown in Figure 2e,f.  Figure 3 shows the electric currents at 0.64 V of the Pt/Ni1−xO/TiN stacks in and HRS states, where Ni1−xO films were deposited at various O2 partial pressu mean values of IHRS and ILRS (red line) increased with the O2 ratio, suggesting Ni1−xO film conductivity depends on the O2 partial pressure, as shown in Figur Ni1−xO films with a 10% O2 fraction required electroforming for resistive switch the Ni1−xO films with a 30% O2 fraction or higher showed reset-first RS behavior electroforming. Figure 3b shows the electrical currents at 0.64 V in the LRS state has a similar tendency to the IHRS with O2 partial pressure, but the slope was lo that of the IHRS state. The IHRS and ILRS showed the highest values for Ni1−xO films d under the 50% O2 partial pressure fraction. To understand the nature of resistance switching, HRS and LRS resistan measured from devices with BE of 0.18, 0.38, 2.00, and 3.69 µm 2 at a bias of ±0.48 V 4a shows the area dependent resistance for BRS device with Ni1−xO films deposite O2 partial pressure fraction. The resistance of the HRS remained almost constant creasing geometric device area, while that of the LRS is almost independent of th area. These area-independent characteristics imply that resistance switching thr device occurs in local regions, such as filament paths, rather than homogeneously uted switching paths [53][54][55][56][57]. Meanwhile, the resistances of reset-first RS devi Ni1−xO films deposited at 50% O2 partial pressure showed increased dependenc device area, as shown in Figure 4b. Because the area dependence of the LRS fo films with 50% O2 partial pressure is close to that of Ni1−xO films with 10% O2 par sure, the nature of the RS is filamentary in the local area. The significant depen HRS on the Ni1−xO films with 50% O2 partial pressure is attributed to the red sistance of the Ni1−xO films, as shown in Figure 4b. The DC, and AC endurance characteristics of the Ni 1−x O device are shown in Figure S1. DC endurance in Figure S1a was measured at a read voltage (V read ) of ±0.25 V under a compliance current of 10 mA. The measured I HRS /I LRS ratio is higher than 10 1 even after 10 3 cycles. Figure S1b shows the AC endurance under pulse, which is measured with a set pulse of −0.95 V with 180 ns, a reset pulse of 1.2 V with 180 ns, and a V read of 0.3 V conditions. The device has a uniform I HRS /I LRS ratio even after 10 5 cycles, which results in a stable RS property. C-AFM measurements investigated the two-dimensional (2D) variation of th film conductivity. Figure 5a illustrates the scheme of the C-AFM measuremen and NiO/SiO2/Pt stacks were simultaneously formed on a sample to compare th ences during the current image mapping. Cross-sectional TEM images of the Ni1 for C-AFM measurements are shown in Figure 5b. The sample-to-sample variati Ni1−xO thickness on the SiO2/Pt stacks was estimated to be within 15%. Therefor nore the difference in conductivity due to thickness variation.  Figure 5d,e. In particular, the current distri relatively uniform in Ni1−xO film with a 50% O2 fraction. In contrast, films deposit 10% O2 partial pressure fraction showed improved resistivity, as shown in Figur C-AFM measurements investigated the two-dimensional (2D) variation of the Ni 1−x O film conductivity. Figure 5a illustrates the scheme of the C-AFM measurement. NiO/Pt and NiO/SiO 2 /Pt stacks were simultaneously formed on a sample to compare the differences during the current image mapping. Cross-sectional TEM images of the Ni 1−x O films for C-AFM measurements are shown in Figure 5b. The sample-to-sample variation in the Ni 1−x O thickness on the SiO 2 /Pt stacks was estimated to be within 15%. Therefore, we ignore the difference in conductivity due to thickness variation.  [58,59].
The proportion of the Ni 3+ state was estimated from the ratio of the Ni 3+ peak area to the Ni 2+ peak area. The Ni 3+ valence state increased while the fraction of Ni 2+ ions decreased with increasing O 2 partial pressure (Figure 6a-c). The Ni 3+ ratio in the film grown under 10% and 50% O 2 partial pressure was estimated at 14.0% and 23.9%, respectively. Meanwhile, the Ni 0 state at the 852.5 eV peak was not observed in our Ni 2p 2/3 peak analysis, although it was considered a conductive path in previous studies [39][40][41]. Conventionally, Ni vacancies form in Ni-deficient NiO films with relatively excessive oxygen. It was reported that nickel deficiency could promote the further oxidation of Ni 2+ ions, which can be expressed with Kröger-Vink notation, as follows [48,49]: where Ni represent Ni 2+ , Ni 3+ , O 2− , and ionized Ni vacancies, respectively. Ni 2+ ions react with oxygen to generate ionized nickel vacancies and two Ni 3+ ions, which affect the conductivity of the nickel oxide films. Therefore, it is shown that the increase Nanomaterials 2022, 12, 2231 6 of 9 in Ni 3+ in Ni 1−x O films is related to the increase in the current in the HRS state of MIM devices and C-AFM. It is expected that Ni deficiency in Ni 1−x O films grown under high O 2 partial pressure causes a high Ni 3+ concentration, leading to a highly conductive state and possibly the reset-first RS behavior with reinforced localized conductive paths [39,60,61]. Further investigation is required to understand how excess Ni 3+ ions produce the reset-first resistive switching behavior in Ni 1−x O films. The effect of the O2 partial pressure on the chemical bonding states in the Ni1− is investigated through XPS analysis. Figure 6a-c show the Ni 2p3/2 peaks of Ni1− deposited with various O2 partial pressures. Ni 0 , Ni 2+ and Ni 3+ states with binding of 852.5, 853.7, and 855.5 eV, respectively, are used for deconvolution of Ni 2p [58,59]. Figure 6. XPS peaks of Ni 2p3/2 of Ni1−xO films with oxygen partial pressure fraction of (a 30% (c) 50%. The effect of the O2 partial pressure on the chemical bonding states in the Ni1−xO is investigated through XPS analysis. Figure 6a-c show the Ni 2p3/2 peaks of Ni1−xO deposited with various O2 partial pressures. Ni 0 , Ni 2+ and Ni 3+ states with binding ene of 852.5, 853.7, and 855.5 eV, respectively, are used for deconvolution of Ni 2p3/2 p [58,59]. The proportion of the Ni 3+ state was estimated from the ratio of the Ni 3+ peak ar the Ni 2+ peak area. The Ni 3+ valence state increased while the fraction of Ni 2+ ion creased with increasing O2 partial pressure (Figure 6a-c). The Ni 3+ ratio in the film g under 10% and 50% O2 partial pressure was estimated at 14.0% and 23.9%, respecti Meanwhile, the Ni 0 state at the 852.5 eV peak was not observed in our Ni 2p2/3 peak ysis, although it was considered a conductive path in previous studies [39][40][41]. Con

Conclusions
In this study, the reset-first RS characteristics of off-stoichiometric Ni 1−x O films were investigated. The RS behavior without the electroforming step was observed in the unipolar and bipolar off-stoichiometric Ni 1−x O films. Ni 3+ distribution contributes significantly to the conductivity of the pristine Ni 1−x O films. The conductivity and Ni deficiency of pristine Ni 1−x O films increased as the O 2 partial pressure increased during a deposition as revealed by the C-AFM and AES results. Moreover, Ni 2+ was further oxidized to Ni 3+ as the O 2 partial pressure increased, as revealed by the XPS results.
The Ni 2 O 3 bonding by Ni 3+ ions is related to the reset-first RS behavior without the electroforming step. This is advantageous in terms of device scale-down, making Ni 1−x O films promising candidates for memory applications by overcoming the limitations of the electroforming step in RRAM.

Conflicts of Interest:
The authors declare no conflict of interest.