Atomic Layer Deposition of Lithium–Nickel–Silicon Oxide Cathode Material for Thin-Film Lithium-Ion Batteries

Lithium nickelate (LiNiO2) and materials based on it are attractive positive electrode materials for lithium-ion batteries, owing to their large capacity. In this paper, the results of atomic layer deposition (ALD) of lithium–nickel–silicon oxide thin films using lithium hexamethyldisilazide (LiHMDS) and bis(cyclopentadienyl) nickel (II) (NiCp2) as precursors and remote oxygen plasma as a counter-reagent are reported. Two approaches were studied: ALD using supercycles and ALD of the multilayered structure of lithium oxide, lithium nickel oxide, and nickel oxides followed by annealing. The prepared films were studied by scanning electron microscopy, spectral ellipsometry, X-ray diffraction, X-ray reflectivity, X-ray photoelectron spectroscopy, time-of-flight secondary ion mass spectrometry, energy-dispersive X-ray spectroscopy, transmission electron microscopy, and selected-area electron diffraction. The pulse ratio of LiHMDS/Ni(Cp)2 precursors in one supercycle ranged from 1/1 to 1/10. Silicon was observed in the deposited films, and after annealing, crystalline Li2SiO3 and Li2Si2O5 were formed at 800 °C. Annealing of the multilayered sample caused the partial formation of LiNiO2. The obtained cathode materials possessed electrochemical activity comparable with the results for other thin-film cathodes.


Introduction
The improvement of small, low-power devices (biosensors [1], smart watches, radio-frequency identification RFID tags, Internet of Things, etc., with power requirements below 10 mW [2]) can be achieved by the development of power sources to provide autonomous operation. Lithium-ion batteries (LIBs), owing to their high energy density, cycle-life, and operational temperature range, are widely applied to power portable electronics. Considering these advantages, LIBs can be regarded as perspective power sources for the abovementioned small-sized devices. The power supply requirements are determined by device construction, functions, and operating conditions. Compact LIBs can be fabricated using traditional electrode manufacturing technology, such as a conventional casting approach [3]. For instance, Wyon produces lithium-ion cells of 6.3 mm 3 with 160 µAh capacity and 94 Wh/L [4] energy density. Smaller LIBs can be manufactured using semiconductor technology.
LIBs with thin-film solid-state construction (TFSSLIBs) have been under development for many years [5]. Some prototypes and products are commercially available but have not yet appeared on the spectra, the thicknesses of the films were calculated. The errors of the film thickness calculation were no more than 0.3 nm. The gradient of the thickness (GT) was calculated using Equation (1): where T max and T min are the maximum and minimum film thicknesses, respectively [16]. X-ray reflectometry (XRR) and x-ray diffraction (XRD) studies were performed on a D8 DISCOVER (Cu-Kα) diffractometer (Bruker, Billerica, MA, USA). Surface-sensitive grazing incidence XRD (GIXRD) modes was used for XRD measurements using 2θ range of 15-65 • with a step of 0.1 • . Exposure time at each step was 1 s. The incidence angle of the primary X-ray beam was 0.7 • . XRR measurements were performed in an angles range 0.3-5 • (increment 0.01 • ) using symmetric scattering geometry. The results of XRD measurements were processed by the Rietveld method using the TOPAS software package (5, Bruker, Billerica, MA, USA)and XRR curves were fitted by the simplex method using LEPTOS (ver. 7.7, Bruker, Billerica, MA, USA).
X-ray photoelectron spectra (XPS) were acquired with a Escalab 250Xi spectrometer (Thermo Fisher Scientific, Waltham, MA, USA). For depth profiling studies, the samples were sputtered by Ar + ions with an energy of 3 keV, for sputtering times ranging from 15 to 1035 s. The samples were excited by Al Kα (1486.7 eV) X-rays in a vacuum no more 7 × 10 −8 Pa.
Elemental depth profiling was also carried out with a time-of-flight secondary ion mass spectrometer (TOF SIMS 5 instrument, ION-TOF GmbH, Münster, Germany). Cs (0.5 keV, area 120 × 120 µm 2 ) and O 2 (0.5 keV, area 150 × 150 µm 2 ) were used for sputtering. The measurements of depth profiles were performed by dynamic SIMS mode using the primary ion gun (Bi + at an energy of 30 keV and a probe measured sample current of 3.1 pA, detection area 100 × 100 µm 2 ).
Scanning electron micrographs of planar and cross-sectional views were obtained using a Merlin scanning electron microscope (SEM, Zeiss, Oerzen, Germany) with a Gemini-II column and a field emission cathode, and a Mira3 SEM (Tescan, Brno, Czech Republic). Both SEMs had field emission cathodes. The SEM spatial resolution was approximately 0.8 nm at an accelerating voltage of 15 kV. A total of 3-4 randomly selected positions on the surface of the sample were investigated. Everhart-Thornley and InLens secondary electron detectors were used for SEM studies. Energy-dispersive X-ray (EDX) analysis was performed using an INCA X-act (Oxford Instruments, High Wycombe, UK) installed on the Zeiss Merlin SEM. A Zeiss Auriga focused ion beam scanning electron microscope (FIB-SEM) dual-beam station was used for lamella preparation for transmission electron microscopy (TEM) studies. The atomic structure, EDX, and electron diffraction patterns were investigated on a Zeiss LIBRA 200FE TEM with an accelerating voltage of 200 kV equipped with an Oxford Instruments INCA X-Max system.
The electrochemical activity of stainless-steel substrates with deposited LNO and LNO-M coatings before and after annealing was studied in coin cells (CR2032). Metallic lithium was used as the counter electrode. Porous polyolefin film (2325, Celgard, Charlotte, NC, USA), and TC-E918 (Tinci, Guangzhou, China) were used as the separator, and electrolyte, respectively. The composition of TC-E918 was a 1-M solution of LiPF 6 in a mixture of organic carbonates. The coin cells were assembled using an OMNI-LAB argon glove box (VAC, Hawthorne, CA, USA); the H 2 O content was less than five ppm. Cyclic voltammetry (CV) was studied using a potentiostat PGSTAT302N+ (Autolab, Utrecht, The Netherlands) in a range of 2.5-4.3 V with 0.5 mV/s scan rate. Charge/discharge cycling was performed using calibrated channels of a CT-3008W-5 V 10 mA battery testing system (Neware, Shenzhen, China) at room temperature, in potential and current ranges of 3.0-4.3 V and 20-80 µA, respectively.

Atomic Layer Deposition and Growth Characteristics
We have deposited pristine lithium (LO) and nickel (NO) oxides, as well as LNO nanolaminates using a supercycle approach. The number of cycles and supercycles, as well as film thicknesses determined by ellipsometry, are given in Table 1. For some selected samples, the film thicknesses were measured by XRR and SEM. The differences among film thickness determined by ellipsometry, SEM, and XRR did not exceed a few percent. According to the SEM images, the deposited films were conformal and uniform, apart from the LNO-1/10 sample, on the surface of which particles of size 10-100 nm were observed (Figures S1 and S2). As the growth per cycle (GPC) of pristine nickel oxide (0.0118 ± 0.0010 nm) is significantly less than that of pristine lithium oxide (0.1225 ± 0.0008 nm), the average GPC of LNO naturally decreases with an increase in the number of Ni(Cp) 2 pulses in one supercycle (Figure 1a). The obtained experimental GPC values are very close to the values calculated based on the GPC of pristine nickel and lithium oxide [59,60].
The growth rates per supercycle (GPSC) naturally increase with the number of Ni(Cp) 2 pulses (Figure 1b), because in this case the total number of precursor pulses in the supercycle increases. The obtained experimental values are also very close to the calculated values.

ALD of LNO Thin Films
3.1.1. Atomic Layer Deposition and Growth Characteristics.
We have deposited pristine lithium (LO) and nickel (NO) oxides, as well as LNO nanolaminates using a supercycle approach. The number of cycles and supercycles, as well as film thicknesses determined by ellipsometry, are given in Table 1. For some selected samples, the film thicknesses were measured by XRR and SEM. The differences among film thickness determined by ellipsometry, SEM, and XRR did not exceed a few percent. According to the SEM images, the deposited films were conformal and uniform, apart from the LNO-1/10 sample, on the surface of which particles of size 10-100 nm were observed (Figures S1 and S2). As the growth per cycle (GPC) of pristine nickel oxide (0.0118 ± 0.0010 nm) is significantly less than that of pristine lithium oxide (0.1225 ± 0.0008 nm), the average GPC of LNO naturally decreases with an increase in the number of Ni(Cp)2 pulses in one supercycle (Figure 1a). The obtained experimental GPC values are very close to the values calculated based on the GPC of pristine nickel and lithium oxide [59,60].
The growth rates per supercycle (GPSC) naturally increase with the number of Ni(Cp)2 pulses (Figure 1b), because in this case the total number of precursor pulses in the supercycle increases. The obtained experimental values are also very close to the calculated values.

Chemical Composition of the Films Determined Using XPS and SEM-EDX
The chemical composition of the film surfaces was studied by XPS. The surface of the films consists mainly of lithium, oxygen, and carbon (Table 2). It should be noted that the LNO samples are Energies 2020, 13, 2345 6 of 24 characterized by a huge concentration of carbon (above 50 at.%). This value significantly exceeds the carbon concentration for the LO samples (about 31 at.%). The lithium concentration was significant and varied in the range of 20-30 at.%, whereas nickel was found only in samples LNO-1/3 and LNO-1/10, and its content did not exceed one percent. Small quantities of silicon (1-2 at.%) and nitrogen (<0.5 at.%) were observed, indicating the presence of impurities derived from lithium precursors (LiHMDS). The peaks of the Li1s spectra ( Figure 2a) are wide and have low intensity. Additionally, the Li 2 O, Li 2 CO 3, and LiOH peak positions are very close to each other [61]. Therefore, it is impossible to carry out the procedure of decomposition and fitting of the spectra to determine the chemical state of lithium with high reliability. However, the positions of the peak maxima suggest that lithium is not in the Li 2 O (55.5 eV) state, but mainly in the LiOH (55.0 eV) and/or Li 2 CO 3 (55.2 eV) state.
Energies 2020, 13, x FOR PEER REVIEW 6 of 25 3.1.2. Chemical composition of the films determined using XPS and SEM-EDX The chemical composition of the film surfaces was studied by XPS. The surface of the films consists mainly of lithium, oxygen, and carbon (Table 2). It should be noted that the LNO samples are characterized by a huge concentration of carbon (above 50 at.%). This value significantly exceeds the carbon concentration for the LO samples (about 31 at.%). The lithium concentration was significant and varied in the range of 20-30 at.%, whereas nickel was found only in samples LNO-1/3 and LNO-1/10, and its content did not exceed one percent. Small quantities of silicon (1-2 at.%) and nitrogen (<0.5 at.%) were observed, indicating the presence of impurities derived from lithium precursors (LiHMDS). The peaks of the Li1s spectra ( Figure 2a) are wide and have low intensity. Additionally, the Li2O, Li2CO3, and LiOH peak positions are very close to each other [61]. Therefore, it is impossible to carry out the procedure of decomposition and fitting of the spectra to determine the chemical state of lithium with high reliability. However, the positions of the peak maxima suggest that lithium is not in the Li2O (55.5 eV) state, but mainly in the LiOH (55.0 eV) and/or Li2CO3 (55.2 eV) state.   The maximum of the O1s spectra is situated in the 531-532 eV range (Figure 2b), and thus it can include the maxima characterizing the states of oxygen in Li 2 CO 3 (532.1 eV) and LiOH (531.2 eV) and does not include the peak Li 2 O at 528.6 eV [62]. Thus, the lithium atoms observed by XPS are a part of Li 2 CO 3 and LiOH. According to the C1s spectra (Figure 2c), the binding energy of most of the carbon (284.8 eV) corresponds to C-C and C-H, i.e., aliphatic hydrocarbons [61]. Moreover, the spectra include peaks of oxygen-containing components: carbonate (290.2 eV) and aldehyde or carboxyl (288.5 eV) [61,62]. The observed carbon is caused by not only the presence of adventitious hydrocarbons [61], but also the moderate content of LiHMDS ligand residues. The assumption is supported by small amounts of silicon and nitrogen found in the deposited films.
According to the peak maxima of the Si2p spectra (Figure 2d), silicon is predominantly in the silicate form but a small amount of silicon may also form silicon dioxide. The N1s peaks ( Figure S3a) are wide and have low intensity; this does not allow an unambiguous determination of the chemical state of nitrogen, which can be bonded with carbon (399 eV) [61] or silicon (398.5 eV) [63], but does not form nitrite or nitrate.
Unfortunately, the peaks of the Ni2p spectra ( Figure S3b) were insufficiently intense to enable drawing any conclusions about the state of nickel in the LNO-1/3 and LNO-1/10.
Because XPS examines only the surface layer of the sample, we studied the bulk composition of the obtained films on the silicon substrate using SEM-EDX ( Figure 3). The electron penetration depth at this energy is approximately 250 nm, and consequently we also observed a signal from the Si substrate. It was found that LNO-1/1 and LNO-1/2 did not comprise any nickel in the bulk of the film. LNO-1/3 and LNO-1/10 contain a small amount of nickel. However, the intensity of the nickel peaks for these samples is several orders of magnitude lower than observed for a multilayer sample LNO-M (to be described later), which includes a nickel oxide layer with a thickness of approximately 23 nm. Probably, the low nickel amount is caused the difficulties of the NiCp 2 chemisorption on the surface species formed after the chemisorption of LiHDMS and exposure to oxygen plasma. Indeed, a number of works indicate that cyclopentadienyl (Cp) ligated precursors which need to break the aromaticity and resonance of the Cp anion to form HCp, often have very long nucleation delays [64]. Probably the nucleation delay leads to the fact that NiCp 2 practically are not chemisorbed during 1 and 2 ALD cycles of NiCp/O 2 plasma, and during 3 and 10 cycles chemisorption is started but nickel content is still very low. The maximum of the O1s spectra is situated in the 531-532 eV range (Figure 2b), and thus it can include the maxima characterizing the states of oxygen in Li2CO3 (532.1 eV) and LiOH (531.2 eV) and does not include the peak Li2O at 528.6 eV [62]. Thus, the lithium atoms observed by XPS are a part of Li2CO3 and LiOH. According to the C1s spectra (Figure 2c), the binding energy of most of the carbon (284.8 eV) corresponds to C-C and C-H, i.e., aliphatic hydrocarbons [61]. Moreover, the spectra include peaks of oxygen-containing components: carbonate (290.2 eV) and aldehyde or carboxyl (288.5 eV) [61,62]. The observed carbon is caused by not only the presence of adventitious hydrocarbons [61], but also the moderate content of LiHMDS ligand residues. The assumption is supported by small amounts of silicon and nitrogen found in the deposited films.
According to the peak maxima of the Si2p spectra (Figure 2d), silicon is predominantly in the silicate form but a small amount of silicon may also form silicon dioxide. The N1s peaks ( Figure S3a) are wide and have low intensity; this does not allow an unambiguous determination of the chemical state of nitrogen, which can be bonded with carbon (399 eV) [61] or silicon (398.5 eV) [63], but does not form nitrite or nitrate.
Unfortunately, the peaks of the Ni2p spectra ( Figure S3b) were insufficiently intense to enable drawing any conclusions about the state of nickel in the LNO-1/3 and LNO-1/10.
Because XPS examines only the surface layer of the sample, we studied the bulk composition of the obtained films on the silicon substrate using SEM-EDX ( Figure 3). The electron penetration depth at this energy is approximately 250 nm, and consequently we also observed a signal from the Si substrate. It was found that LNO-1/1 and LNO-1/2 did not comprise any nickel in the bulk of the film. LNO-1/3 and LNO-1/10 contain a small amount of nickel. However, the intensity of the nickel peaks for these samples is several orders of magnitude lower than observed for a multilayer sample LNO-M (to be described later), which includes a nickel oxide layer with a thickness of approximately 23 nm. Probably, the low nickel amount is caused the difficulties of the NiCp2 chemisorption on the surface species formed after the chemisorption of LiHDMS and exposure to oxygen plasma. Indeed, a number of works indicate that cyclopentadienyl (Cp) ligated precursors which need to break the aromaticity and resonance of the Cp anion to form HCp, often have very long nucleation delays [64]. Probably the nucleation delay leads to the fact that NiCp2 practically are not chemisorbed during 1 and 2 ALD cycles of NiCp/O2 plasma, and during 3 and 10 cycles chemisorption is started but nickel content is still very low. The as-prepared NO and nickel-containing LNO samples (LNO-1/3 and LNO-1/10) were studied by XRD. The GIXRD pattern of NO clearly shows reflections (111), (200), (220), corresponding to cubic Fm-3m modifications of NiO ( Figure S4). However, the LNO samples were amorphous. NiO, Ni2O3, Fm-3m modifications of NiO ( Figure S4). However, the LNO samples were amorphous. NiO, Ni 2 O 3 , Ni 3 O 4 , and LiNiO 2 reflections were not observed. After a rapid annealing process (15 min in air at 800 • C), only reflections of Li 2 SiO 3 ((220), (201), (020) and (021)) and Li 2 Si 2 O 5 ((130), (040), (111), and (002)) were manifested for samples ( Figure 4). Reflections in the region of 26-30 • , 33 • , 48 • , 55 • , 57 • , and 62 • are associated with imperfections in the structure of the silicon. Thus, the nickel in the LNO samples is either in an amorphous state, or its concentration is insufficient for phase detection.
Energies 2020, 13, x FOR PEER REVIEW 8 of 25 Ni3O4, and LiNiO2 reflections were not observed. After a rapid annealing process (15 min in air at 800 °C), only reflections of Li2SiO3 ((220), (201), (020) and (021)) and Li2Si2O5 ((130), (040), (111), and (002)) were manifested for samples ( Figure 4). Reflections in the region of 26-30°, 33°, 48°, 55°, 57°, and 62° are associated with imperfections in the structure of the silicon. Thus, the nickel in the LNO samples is either in an amorphous state, or its concentration is insufficient for phase detection.  (Table 3). Both layers have a low roughness (1.2-2.4 nm). Nevertheless, the upper layer has a small thickness (5-6 nm) and a remarkable density gradient (top: 1.7-2.3 g/cm 3 , bottom: 0.5-1.2 g/cm 3 ). The bottom layers are thicker (39 and 80 nm) and denser (2.5-3 g/cm 3 ). The densities of the inner layers are higher than the densities of bulk Li2O (2.01 g/cm 3 ), Li2CO3 (2.11 g/cm 3 ), and LiOH (1.45 g/cm 3 ), but much less than those of NiO (6.67 g/cm 3 ), Ni2O3 (5.18 g/cm 3 ), and LiNiO2 (4.81 g/cm 3 ), and are close to that of Li2SiO3 (2.52 g/cm 3 ). It is probable that the upper layers are a mixture of lithium carbonates, and the bottom layers are a mixture of lithium oxide, silicates, and a small amount of nickel oxides and/or lithium nickelates. After annealing at 800 °C, the samples have one homogeneous layer without a significant density gradient. It is clear that the upper layer of carbonates is decomposed during annealing. The density of the remaining layer increases. For example, the density range of LNO-1/3 increased from 2.46-2.49 to 2.57-2.78 g/cm 3 .  (Table 3). Both layers have a low roughness (1.2-2.4 nm). Nevertheless, the upper layer has a small thickness (5-6 nm) and a remarkable density gradient (top: 1.7-2.3 g/cm 3 , bottom: 0.5-1.2 g/cm 3 ). The bottom layers are thicker (39 and 80 nm) and denser (2.5-3 g/cm 3 ). The densities of the inner layers are higher than the densities of bulk Li 2 O (2.01 g/cm 3 ), Li 2 CO 3 (2.11 g/cm 3 ), and LiOH (1.45 g/cm 3 ), but much less than those of NiO (6.67 g/cm 3 ), Ni 2 O 3 (5.18 g/cm 3 ), and LiNiO 2 (4.81 g/cm 3 ), and are close to that of Li 2 SiO 3 (2.52 g/cm 3 ). It is probable that the upper layers are a mixture of lithium carbonates, and the bottom layers are a mixture of lithium oxide, silicates, and a small amount of nickel oxides and/or lithium nickelates. After annealing at 800 • C, the samples have one homogeneous layer without a significant density gradient. It is clear that the upper layer of carbonates is decomposed during annealing. The density of the remaining layer increases. For example, the density range of LNO-1/3 increased from 2.46-2.49 to 2.57-2.78 g/cm 3 .

ALD and Growth Characteristics of Multilayered LNO Thin Films
As the use of the supercycle approach for ALD of lithium nickelate was unsuccessful, we tried to prepare the multilayer structures of nickel and lithium oxides (Table 1) with subsequent annealing.
The sample prepared according to this scheme will hereinafter be referred to as LNO-M.

ALD and growth characteristics of multilayered LNO thin films
As the use of the supercycle approach for ALD of lithium nickelate was unsuccessful, we tried to prepare the multilayer structures of nickel and lithium oxides (Table 1) with subsequent annealing.
The sample prepared according to this scheme will hereinafter be referred to as LNO-M.

ALD and growth characteristics of multilayered LNO thin films
As the use of the supercycle approach for ALD of lithium nickelate was unsuccessful, we tried to prepare the multilayer structures of nickel and lithium oxides (Table 1) with subsequent annealing.
The sample prepared according to this scheme will hereinafter be referred to as LNO-M.   at 51 • , as well as the peak at 38 • , which appear upon annealing, could fit the reflections of Ni 2 O 3 , but there are no other intense peaks of Ni 2 O 3 at 32 • , 45 • , and 57 • . Therefore, we believe that the peaks at 51 • and 53 • correspond to defects in the silicon substrate. After annealing, the intensity of the (111) NiO peak increased, and several peaks appeared on the pattern. The small peaks at 22-23 • and the peak at 38 • probably correspond to lithium silicide, Li 22

Spectral Ellipsometry and X-ray Reflectometry
Based on the GPC of the LO, NO, and LNO-1/3 samples, one could expect that the thickness of the obtained LNO-M would be approximately 100.9 nm (calculated thickness, Figure 5). However, the total film thickness obtained by modeling the ellipsometry spectrum and fitting was 88.4 nm. According to reflectometry data fitting ( Figure S7, Table 4), the total film thickness is approximately 83.8 nm ( Figure 5, green), which correlates with the ellipsometry results, but significantly less than the expected value of 100.9 nm. This mismatch can be explained by the difference in the nucleation effect on different substrates. Because of this, the growth rates of LO and LNO-1/3 layers on the surface of already deposited NO are much lower than those on the silicon surface. Similarly, the growth of the upper layer of NO on the surface of LNO-1/3 can also be slowed down owing to the nucleation effect. This assumption is confirmed by the data obtained from the spectral ellipsometry using witnesses (Figure 7), which we used during the ALD of LNO-M. The film thickness without the bottom layer of NO was 72 nm, and the thickness of the film with only the NO layer was 33.3 nm. These values are significantly higher than the thicknesses determined for the LNO-M sample, i.e., 60.5 and 27.9 nm, respectively ( Figure 5).
Energies 2020, 13, x FOR PEER REVIEW 10 of 25 The cell parameters calculated by the Rietveld method were 0.4174 and 0.4185 nm for LNO-M and LNO-M-800, respectively. These values are very close to the cell parameter for the NO sample (0.4168 nm). The diffraction patterns of the as-deposited sample revealed peaks at 51° and 53°. The peak at 51°, as well as the peak at 38°, which appear upon annealing, could fit the reflections of Ni2O3, but there are no other intense peaks of Ni2O3 at 32°, 45°, and 57°. Therefore, we believe that the peaks at 51° and 53° correspond to defects in the silicon substrate. After annealing, the intensity of the (111) NiO peak increased, and several peaks appeared on the pattern. The small peaks at 22-23° and the peak at 38° probably correspond to lithium silicide, Li22Si5. Three peaks located at 23-25° are related to the most intense reflections (130), (040), and (111) of Li2Si2O5.

Spectral ellipsometry and X-ray reflectometry
Based on the GPC of the LO, NO, and LNO-1/3 samples, one could expect that the thickness of the obtained LNO-M would be approximately 100.9 nm (calculated thickness, Figure 5). However, the total film thickness obtained by modeling the ellipsometry spectrum and fitting was 88.4 nm. According to reflectometry data fitting ( Figure S7, Table 4), the total film thickness is approximately 83.8 nm ( Figure 5, green), which correlates with the ellipsometry results, but significantly less than the expected value of 100.9 nm. This mismatch can be explained by the difference in the nucleation effect on different substrates. Because of this, the growth rates of LO and LNO-1/3 layers on the surface of already deposited NO are much lower than those on the silicon surface. Similarly, the growth of the upper layer of NO on the surface of LNO-1/3 can also be slowed down owing to the nucleation effect. This assumption is confirmed by the data obtained from the spectral ellipsometry using witnesses (Figure 7), which we used during the ALD of LNO-M. The film thickness without the bottom layer of NO was 72 nm, and the thickness of the film with only the NO layer was 33.3 nm. These values are significantly higher than the thicknesses determined for the LNO-M sample, i.e., 60.5 and 27.9 nm, respectively ( Figure 5). No clear boundaries were detected between the LO and LNO-1/3 layers according to ellipsometry and XRR data fitting ( Figure S7). Three layers are presented in the bulk of the film, which differ in density ( Table 4). The bottom layer (layer 1) has a density of 6.0-6.6 g/m 3 , which is close to that of NiO (6.67 g/cm 3 ). Presumably, the middle layer (density 1.86-2.31 g/m 3 ) contains lithium oxides, hydroxides, silicates, and a small amount of nickel; the denser (2.55 g/cm 3 ) outer layer consists of lithium carbonate.  No clear boundaries were detected between the LO and LNO-1/3 layers according to ellipsometry and XRR data fitting ( Figure S7). Three layers are presented in the bulk of the film, which differ in density ( Table 4). The bottom layer (layer 1) has a density of 6.0-6.6 g/m 3 , which is close to that of NiO (6.67 g/cm 3 ). Presumably, the middle layer (density 1.86-2.31 g/m 3 ) contains lithium oxides, hydroxides, silicates, and a small amount of nickel; the denser (2.55 g/cm 3 ) outer layer consists of lithium carbonate.

Chemical Composition of the Films. XPS and TOF-SIMS Depth Profiling
According to the EDX data, the LNO-M sample contains nickel, the concentration of which is significantly higher than those in the LNO samples (Figure 3). For a more detailed study of the composition of the LNO-M sample, depth profiling was performed by XPS ( Figure 8a) and TOF-SIMS (Figure 8b and Figure S8). The results of TOF-SIMS profiling are similar to those of XPS in terms of the dynamics of ion changes with changes in sputtering time. However, for TOF-SIMS, the boundaries between the layers are not very clear, and the ion yield values strongly depend on the nature of the element and sputtering ions, which complicates the quantitative analysis. A qualitative elemental analysis was performed based on the XPS data (Tables S1 and S2).

Chemical composition of the films. XPS and TOF-SIMS depth profiling
According to the EDX data, the LNO-M sample contains nickel, the concentration of which is significantly higher than those in the LNO samples (Figure 3). For a more detailed study of the composition of the LNO-M sample, depth profiling was performed by XPS ( Figure 8a) and TOF-SIMS (Figures 8b, S8). The results of TOF-SIMS profiling are similar to those of XPS in terms of the dynamics of ion changes with changes in sputtering time. However, for TOF-SIMS, the boundaries between the layers are not very clear, and the ion yield values strongly depend on the nature of the element and sputtering ions, which complicates the quantitative analysis. A qualitative elemental analysis was performed based on the XPS data (Tables S1 and S2). The upper layer of the LNO-M should consist of nickel oxide ( Figure 5). However, the XPS data did not show the presence of this element in the surface layer (Table S1, Figure 8a). The surface contains a high concentration of carbon, which decreases rapidly with increasing depth into the film. Carbon was not detected at sputtering times of 110 s and longer. A significant concentration of carbonate ions was also recorded on the surface according to the TOF-SIMS ( Figure S8 LiCO3 − ions, black open circles). Obviously, the source of this carbon is not only surface contamination, but also lithium carbonization during its exposure to the air. We assume that the carbonization led to the absence of nickel on the surface. The upper NO layer was grown, but it was not continuous and consisted of particles. As a result, these NO particles were overgrown (coated) with a layer of lithium carbonate and were no longer on the surface, but in the bulk of the film. In addition to decreasing the carbon as the sputtering time increases, the lithium concentration decreases with increasing depth in the film, whereas the concentrations of oxygen and silicon increase. There is a transition from carbonate to lithium and silicon oxides/hydroxides. Apparently, this layer corresponds to the surface layer (layer 3) detected by XRR (Table 4).
Starting from 110 s of XPS sputtering, the next layer (layer 2 at Table 4) is reached, and the reverse trend of concentration changes is observed during sputtering, i.e., gradual decreases in the amounts of oxygen and silicon and an increase in the lithium are seen. The concentration of the silicon in this bulk layer is very high (21-24%). Obviously, the source of the silicon is unreacted -Si(CH3) ligands of LiHMDS. Previously was shown [65] that with the use of LiHMDS and ozone for ALD, the partial formation of SiO2 is possible. In work [66], it was shown the possibility of production Li2SiO3 films using LiHMDS and oxygen source. Thereby the bulk of LNO samples also contain a large amount of silicon, but only a small amount (1-2 at.%) on the surface (Table 2). Therefore, the formation of crystalline phases Li2SiO3 and Li2Si2O5 (Figure 4b) for LNO and LNO-M samples after annealing is quite expected. The nickel concentration in this layer does not exceed 0.1-0.2% (Table S1). According to both the TOF-SIMS and XPS, there are no clear boundaries between the LO and LNO-1/3, although the TOF-SIMS data show a significant change in the trend of Li + and Li2O + ion concentration near 11 the source of this carbon is not only surface contamination, but also lithium carbonization during its exposure to the air. We assume that the carbonization led to the absence of nickel on the surface. The upper NO layer was grown, but it was not continuous and consisted of particles. As a result, these NO particles were overgrown (coated) with a layer of lithium carbonate and were no longer on the surface, but in the bulk of the film. In addition to decreasing the carbon as the sputtering time increases, the lithium concentration decreases with increasing depth in the film, whereas the concentrations of oxygen and silicon increase. There is a transition from carbonate to lithium and silicon oxides/hydroxides. Apparently, this layer corresponds to the surface layer (layer 3) detected by XRR (Table 4). Starting from 110 s of XPS sputtering, the next layer (layer 2 at Table 4) is reached, and the reverse trend of concentration changes is observed during sputtering, i.e., gradual decreases in the amounts of oxygen and silicon and an increase in the lithium are seen. The concentration of the silicon in this bulk layer is very high (21-24%). Obviously, the source of the silicon is unreacted -Si(CH 3 ) ligands of LiHMDS. Previously was shown [65] that with the use of LiHMDS and ozone for ALD, the partial formation of SiO 2 is possible. In work [66], it was shown the possibility of production Li 2 SiO 3 films using LiHMDS and oxygen source. Thereby the bulk of LNO samples also contain a large amount of silicon, but only a small amount (1-2 at.%) on the surface (Table 2). Therefore, the formation of crystalline phases Li 2 SiO 3 and Li 2 Si 2 O 5 (Figure 4(2)) for LNO and LNO-M samples after annealing is quite expected. The nickel concentration in this layer does not exceed 0.1-0.2% (Table S1). According to both the TOF-SIMS and XPS, there are no clear boundaries between the LO and LNO-1/3, although the TOF-SIMS data show a significant change in the trend of Li + and Li 2 O + ion concentration near 11 s of sputtering ( Figure 8b). However, these areas could be TOF-SIMS artifacts, as a decrease in the number of ions is observed for all positive ions, and for negative, we do not see a similar trend.
After 675 s of sputtering during the XPS measurements, the nickel concentration increases sharply, and the silicon and lithium have disappeared. The Ni/O ratio becomes close to 1. Obviously, we are approaching the bottom layer (layer 1 in Table 4), which consists of nickel oxide. The selected sputtering time did not allow reaching the substrate during the XPS study.
In the TOF-SIMS depth profiles (Figure 8b and Figure S8), the transition from the nickel oxide layer to the stainless-steel substrate is visible. However, a clear boundary is not observed; therefore, the boundaries between the layers in Figure 8b and Figure S8 can be considered as approximated. In the transition region from nickel oxide to SS316, a gradual decrease in the concentrations of Ni + , NiO + , Li 2 O + , NiO − , SiO 2 − , and LiO − ions is observed. The number of FeO − ions augmented and reaches a maximum, i.e., we reach a layer of iron oxide that is always present on the steel surface. Further sputtering is accompanied by a decrement in FeO − and an increase in Fe + , and thus we have reached the steel substrate. It is noteworthy that a sharp increase in the number of Li + ions is also observed in this transition layer. This may be caused by the diffusion of Li + into the substrate or at its surface through the nickel oxide. Indeed, it was shown in [67] that in ALD-deposited films of Li 2 CO 3 , lithium can diffuse through thin layers of HfO 2 and ZrO 2 after annealing. Moreover, a number of works [46,47] showed that diffusion can occur even, at relatively low temperatures (225 • C), directly during the ALD process. Thus, sequential ALD using Li(thd) and O 3 or LiO t Bu and H 2 O on previously deposited MnO 2 or V 2 O 5 films resulted in postlithiation of the films rather than an additional layer of Li 2 CO 3 .
For an annealed sample (LNO-M-800), the composition practically does not change with respect to depth (Figures S9 and S10). There are no clear boundaries between the layers, either according to XPS or TOF-SIMS. Thus, upon annealing, mutual diffusion of the elements occurred. A slight difference is observed only for the surface and bulk layers. Oxygen (11.7-13%), lithium (69-73.4%), silicon (0.3-0.8%), nickel (2.4-3.9%), and iron (7-7.8%) (Table S2) are evenly distributed throughout the depth. In this case, the chromium concentration increases with depth from 0.9 to 2.8%. Carbon is mainly present on the surface, but its concentration is much lower than that before annealing. The presence of iron and chromium in the film is obviously caused by the diffusion of these elements from the steel substrate. The huge amount of lithium is obviously overstated. In fact, the Li1s and Fe3p levels are overlapped, and their reliable separation is difficult.

Chemical Composition of the Films. Detailed Study of XPS Spectra
For a more detailed study of the composition of the samples during profiling, the corresponding spectra of the C1s, O1s, Li1s, Ni2p, Si2p, Fe2p, and Cr2p levels were analyzed (Figure 9 and Figure S11).
Energies 2020, 13, x FOR PEER REVIEW 12 of 25 s of sputtering (Figure 8b). However, these areas could be TOF-SIMS artifacts, as a decrease in the number of ions is observed for all positive ions, and for negative, we do not see a similar trend. After 675 s of sputtering during the XPS measurements, the nickel concentration increases sharply, and the silicon and lithium have disappeared. The Ni/O ratio becomes close to 1. Obviously, we are approaching the bottom layer (layer 1 in Table 4), which consists of nickel oxide. The selected sputtering time did not allow reaching the substrate during the XPS study.
In the TOF-SIMS depth profiles (Figures 8b, S8), the transition from the nickel oxide layer to the stainless-steel substrate is visible. However, a clear boundary is not observed; therefore, the boundaries between the layers in Figures 8b, S8 can be considered as approximated. In the transition region from nickel oxide to SS316, a gradual decrease in the concentrations of Ni + , NiO + , Li2O + , NiO − , SiO2 − , and LiO − ions is observed. The number of FeO − ions augmented and reaches a maximum, i.e., we reach a layer of iron oxide that is always present on the steel surface. Further sputtering is accompanied by a decrement in FeO − and an increase in Fe + , and thus we have reached the steel substrate. It is noteworthy that a sharp increase in the number of Li + ions is also observed in this transition layer. This may be caused by the diffusion of Li + into the substrate or at its surface through the nickel oxide. Indeed, it was shown in [67] that in ALD-deposited films of Li2CO3, lithium can diffuse through thin layers of HfO2 and ZrO2 after annealing. Moreover, a number of works [46,47] showed that diffusion can occur even, at relatively low temperatures (225 °C), directly during the ALD process. Thus, sequential ALD using Li(thd) and O3 or LiO t Bu and H2O on previously deposited MnO2 or V2O5 films resulted in postlithiation of the films rather than an additional layer of Li2CO3.
For an annealed sample (LNO-M-800), the composition practically does not change with respect to depth (Figures S9, S10). There are no clear boundaries between the layers, either according to XPS or TOF-SIMS. Thus, upon annealing, mutual diffusion of the elements occurred. A slight difference is observed only for the surface and bulk layers. Oxygen (11.7-13%), lithium (69-73.4%), silicon (0.3-0.8%), nickel (2.4-3.9%), and iron (7-7.8%) (Table S2) are evenly distributed throughout the depth. In this case, the chromium concentration increases with depth from 0.9 to 2.8%. Carbon is mainly present on the surface, but its concentration is much lower than that before annealing. The presence of iron and chromium in the film is obviously caused by the diffusion of these elements from the steel substrate. The huge amount of lithium is obviously overstated. In fact, the Li1s and Fe3p levels are overlapped, and their reliable separation is difficult.

Chemical composition of the films. Detailed study of XPS spectra
For a more detailed study of the composition of the samples during profiling, the corresponding spectra of the C1s, O1s, Li1s, Ni2p, Si2p, Fe2p, and Cr2p levels were analyzed (Figures 9 and S11).   carbonate (maximum at 532.1 eV). Then (45 s sputtering), the peak broadens owing to the manifestation of two or even three components: silicate/silica (532.8 eV) [68,69], lithium carbonate (532.1 eV) and probably lithium hydroxide (531.1 eV). For the longer sputtering times, the silicate/silica component becomes the most intense, but lithium hydroxide is also present. With an increase in the etching time to 675 s, a shift to the lithium hydroxide is observed, and at 1035 s, an intense peak of NiO appears (529.6 eV) [68,69].
For the annealed sample (LNO-M-800) (Figure 9d), oxygen on the surface (without sputtering) is predominantly in the form of the carbonate, but the intensity of this peak is low. The spectra measured after different sputtering times are similar to each other. Sputtering leads to increased intensities and peak maxima shifted toward lower energies. These peaks are most likely a combination of NiO, FeO, and LiOH peaks and indicate the presence of a mixture of these oxides and this hydroxide.
Unfortunately, the analysis of the Li1s spectra (Figure 9e,f) is difficult, as the positions of the maxima for Li 2 O, LiOH, and Li 2 CO 3 are very close. For the LNO-M-800 sample, iron appears in the bulk of the film, whose 3p level overlaps with Li1s.
For LNO-M, nickel only appears at 675 s of sputtering (Figure 9f). At 1035 s, the peaks become very intense. For the LNO-M-800 sample (Figure 9g), nickel is visible over the entire depth, and the peaks are intense. However, their exact deconvolution and fitting are difficult, owing to the presence of shake-up and plasmon loss structures, and multiplet splitting [70]. Nevertheless, a qualitative analysis of the spectra leads us to conclude that nickel is present both in the oxide form and in the form of metallic or carbon-bound nickel. For LNO-M, the oxide phase is predominant, and after calcination, the proportion of Ni and/or NI-C increases, probably owing to the diffusion of nickel from the stainless-steel substrate.
For LNO-M, a very high concentration of silicon in the bulk of the film is observed. In the near-surface region ( Figure S11a), silicon is in the form of silicate (maximum near 101.5 eV). However, with increasing depth into the film thickness, the maximum shifts toward higher energies, which are more typical for SiO 2 . For LNO-M-800, silicon is visible only on the surface ( Figure S11b) and in small quantities. As will be shown below, silicon diffuses deep into the film.
For LNO-M-800, the XPS spectra of the Fe2p and Cr2p levels also show iron and chromium in an oxide state ( Figure S11c,d); thus, these elements also diffuse to the deposited film during the annealing process.

SEM and TEM. Morphology, Local Structure and Composition of LNO-M-800 Thin Film
LNO-M-800 deposited on an SS316 support was selected for detailed study by SEM, TEM, and selected-area electron diffraction (SAED). A SEM image (Figure 10a) shows that the SS316 surface is inhomogeneously coated. There are two types of structures: (1) crystals of size 50-250 nm (see inset 1 in Figure 10a), which form a loose surface, and (2) triangular crystals lateral size 250-500 nm (see inset 2 in Figure 10a). The latter form a stepped surface and their agglomerates spread over a few tens of microns. A probable reason for these two surface types may be the grain structure of the steel substrate leading to local inhomogeneity of the coatings [71]. The stepped surface (inset 2 in Figure 10a) was chosen for lamella preparation, because it has large crystallites with a perfect lattice structure. Detailed SEM images of the lamella preparation process performed by a Zeiss Auriga dual-beam station are presented in Figure S12. TEM was used in two modes: conventional transmission mode to obtain images with highresolution and SAED studies and scanning transmission mode for EDX mapping. A bright-field (BF) scanning transmission electron microscopy (STEM) image of the lamella is demonstrated in Figure  10b. Two inset images show higher magnification. In the left inset, three typical layers are marked. The similar layers are also clearly visible on right insert in Figure 10b and throughout the lamella. The first layer is the stainless-steel substrate with typical polycrystals of metal that have black-white contrast. The second layer is amorphous with thickness of 50-80 nm and uniform gray contrast, their origin will be discussed below. The third layer is the layer formed by ALD and subsequent annealing. Crystals of the third layer have a height from 50 to 200 nm. Their lateral sizes, shapes, and orientations show variation. Figure 11a presents a high-angle annular dark-field (HAADF) STEM image of the area used for EDX mapping. EDX maps of Ni, O, Fe, Cr, Si are shown in Figure 11b-f, respectively. The uniform distributions of Ni and O in the crystal volume correspond with the XPS and SIMS depth profiling results. The amorphous layer (Figure 10b, layer 2) predominantly contains O, Si, and Cr. We proposed that the high concentrations of oxygen and chromium may be caused by corrosion or/and passivation of stainless steel described elsewhere [72]. The silicon content is most likely due to the diffusion from the ALD-deposited film. This diffusion explains the significant decrease in the silicon content after annealing according to XPS depth profiling (Tables S1 and S2). In other words, we did not reach the bottom silicon-containing layer during 600 seconds of XPS sputtering. The Fe and Cr maps demonstrate that these atoms diffused from the substrate to the ALD-deposited layers during annealing, which correlates with the XPS and TOF-SIMS results. Areas 1 and 2 marked in Figure 11a are regions of the EDX spectrum measurements presented in Figure S13. TEM was used in two modes: conventional transmission mode to obtain images with high-resolution and SAED studies and scanning transmission mode for EDX mapping. A bright-field (BF) scanning transmission electron microscopy (STEM) image of the lamella is demonstrated in Figure 10b. Two inset images show higher magnification. In the left inset, three typical layers are marked. The similar layers are also clearly visible on right insert in Figure 10b and throughout the lamella. The first layer is the stainless-steel substrate with typical polycrystals of metal that have black-white contrast. The second layer is amorphous with thickness of 50-80 nm and uniform gray contrast, their origin will be discussed below. The third layer is the layer formed by ALD and subsequent annealing. Crystals of the third layer have a height from 50 to 200 nm. Their lateral sizes, shapes, and orientations show variation. Figure 11a presents a high-angle annular dark-field (HAADF) STEM image of the area used for EDX mapping. EDX maps of Ni, O, Fe, Cr, Si are shown in Figure 11b-f, respectively. The uniform distributions of Ni and O in the crystal volume correspond with the XPS and SIMS depth profiling results. The amorphous layer (Figure 10b, layer 2) predominantly contains O, Si, and Cr. We proposed that the high concentrations of oxygen and chromium may be caused by corrosion or/and passivation of stainless steel described elsewhere [72]. The silicon content is most likely due to the diffusion from the ALD-deposited film. This diffusion explains the significant decrease in the silicon content after annealing according to XPS depth profiling (Tables S1 and S2). In other words, we did not reach the bottom silicon-containing layer during 600 seconds of XPS sputtering. The Fe and Cr maps demonstrate that these atoms diffused from the substrate to the ALD-deposited layers during annealing, which correlates with the XPS and TOF-SIMS results. Areas 1 and 2 marked in Figure 11a are regions of the EDX spectrum measurements presented in Figure S13. TEM was used in two modes: conventional transmission mode to obtain images with highresolution and SAED studies and scanning transmission mode for EDX mapping. A bright-field (BF) scanning transmission electron microscopy (STEM) image of the lamella is demonstrated in Figure  10b. Two inset images show higher magnification. In the left inset, three typical layers are marked. The similar layers are also clearly visible on right insert in Figure 10b and throughout the lamella. The first layer is the stainless-steel substrate with typical polycrystals of metal that have black-white contrast. The second layer is amorphous with thickness of 50-80 nm and uniform gray contrast, their origin will be discussed below. The third layer is the layer formed by ALD and subsequent annealing. Crystals of the third layer have a height from 50 to 200 nm. Their lateral sizes, shapes, and orientations show variation. Figure 11a presents a high-angle annular dark-field (HAADF) STEM image of the area used for EDX mapping. EDX maps of Ni, O, Fe, Cr, Si are shown in Figure 11b-f, respectively. The uniform distributions of Ni and O in the crystal volume correspond with the XPS and SIMS depth profiling results. The amorphous layer (Figure 10b, layer 2) predominantly contains O, Si, and Cr. We proposed that the high concentrations of oxygen and chromium may be caused by corrosion or/and passivation of stainless steel described elsewhere [72]. The silicon content is most likely due to the diffusion from the ALD-deposited film. This diffusion explains the significant decrease in the silicon content after annealing according to XPS depth profiling (Tables S1 and S2). In other words, we did not reach the bottom silicon-containing layer during 600 seconds of XPS sputtering. The Fe and Cr maps demonstrate that these atoms diffused from the substrate to the ALD-deposited layers during annealing, which correlates with the XPS and TOF-SIMS results. Areas 1 and 2 marked in Figure 11a are regions of the EDX spectrum measurements presented in Figure S13. A high-resolution TEM (HR-TEM) image obtained from the region indicated by the black arrow in the right inset of Figure 10b is presented in Figure 12. The original HR-TEM image is difficult to analyze, owing to the non-ideal surface and high thickness. Only weak diagonal periodic lines are clearly observed. A Fourier filtration image of the region marked in the original image is shown in the inset of Figure 12a. In this case, a layering crystal structure without defects is clearly observed. SAED with aperture diameter 80 nm was used to accurately calculate the lattice parameters. The SAED pattern obtained for the same region of the HRTEM image is presented in Figure 12b. The three sets of marked reflexes and calculated interplanar spaces exactly match planes (003), (101), and (104) with distances 4.78, 2.47, and 2.06 Å, respectively, of LiNiO2 (Table S3). A fourth arrow also could be associated with the (10−8) or (2−10) set of LiNiO2 planes. The layering structures were not observed in all HR TEM images and SAED could not help to identify certain phases. Moreover, the presence of Fe and Cr elements in ALD layers (see Figure 11d,e) should distort the lattice parameters and complicate the phase identification process. Considering the XRD results ( Figure 6), which identified a NiO phase in the LNO-M sample, reflexes and corresponding interplanar space in the SAED patterns could also be associated with the NiO crystal lattice parameters (Table S3). However, the (003) set of reflections clearly belongs only to LiNiO2. Comparing the results of XRD (integral method) and SAED (local method), we assume that the deposited material predominantly consists of the NiO phase with inclusions of LiNiO2 crystallites.  A high-resolution TEM (HR-TEM) image obtained from the region indicated by the black arrow in the right inset of Figure 10b is presented in Figure 12. The original HR-TEM image is difficult to analyze, owing to the non-ideal surface and high thickness. Only weak diagonal periodic lines are clearly observed. A Fourier filtration image of the region marked in the original image is shown in the inset of Figure 12a. In this case, a layering crystal structure without defects is clearly observed. SAED with aperture diameter 80 nm was used to accurately calculate the lattice parameters. The SAED pattern obtained for the same region of the HRTEM image is presented in Figure 12b. The three sets of marked reflexes and calculated interplanar spaces exactly match planes (003), (101), and (104) with distances 4.78, 2.47, and 2.06 Å, respectively, of LiNiO 2 (Table S3). A fourth arrow also could be associated with the (10−8) or (2−10) set of LiNiO 2 planes. The layering structures were not observed in all HR TEM images and SAED could not help to identify certain phases. Moreover, the presence of Fe and Cr elements in ALD layers (see Figure 11d,e) should distort the lattice parameters and complicate the phase identification process. Considering the XRD results ( Figure 6), which identified a NiO phase in the LNO-M sample, reflexes and corresponding interplanar space in the SAED patterns could also be associated with the NiO crystal lattice parameters (Table S3). However, the (003) set of reflections clearly belongs only to LiNiO 2 . Comparing the results of XRD (integral method) and SAED (local method), we assume that the deposited material predominantly consists of the NiO phase with inclusions of LiNiO 2 crystallites.  A high-resolution TEM (HR-TEM) image obtained from the region indicated by the black arrow in the right inset of Figure 10b is presented in Figure 12. The original HR-TEM image is difficult to analyze, owing to the non-ideal surface and high thickness. Only weak diagonal periodic lines are clearly observed. A Fourier filtration image of the region marked in the original image is shown in the inset of Figure 12a. In this case, a layering crystal structure without defects is clearly observed. SAED with aperture diameter 80 nm was used to accurately calculate the lattice parameters. The SAED pattern obtained for the same region of the HRTEM image is presented in Figure 12b. The three sets of marked reflexes and calculated interplanar spaces exactly match planes (003), (101), and (104) with distances 4.78, 2.47, and 2.06 Å, respectively, of LiNiO2 (Table S3). A fourth arrow also could be associated with the (10−8) or (2−10) set of LiNiO2 planes. The layering structures were not observed in all HR TEM images and SAED could not help to identify certain phases. Moreover, the presence of Fe and Cr elements in ALD layers (see Figure 11d,e) should distort the lattice parameters and complicate the phase identification process. Considering the XRD results ( Figure 6), which identified a NiO phase in the LNO-M sample, reflexes and corresponding interplanar space in the SAED patterns could also be associated with the NiO crystal lattice parameters (Table S3). However, the (003) set of reflections clearly belongs only to LiNiO2. Comparing the results of XRD (integral method) and SAED (local method), we assume that the deposited material predominantly consists of the NiO phase with inclusions of LiNiO2 crystallites.

Electrochemistry
ALD films deposited on SS316 supports were used to study the electrochemical activity. The ALD samples and steel substrate were investigated with the use of CV. The areas under the anode and cathode curves for the SS316 support and NO samples were almost equal, which indicates good reversibility of underlying electrochemical processes (Figure 13a). The maximum currents observed on the anode and cathode curves of the steel support are 1.5 µA (3.5 V) and 2.0 µA (3.0 V), respectively. According to [73], more detailed CV peaks can be observed using different scan rates. The annealed LNO-M samples showed augmented areas under the anode and cathode curves in comparison with the initial steel and NiO. The positions of the current maxima of the anode and cathode curves shift to 4.3 V (20-26 µA) and 3.75 V (5-10 µA), respectively.

Electrochemistry
ALD films deposited on SS316 supports were used to study the electrochemical activity. The ALD samples and steel substrate were investigated with the use of CV. The areas under the anode and cathode curves for the SS316 support and NO samples were almost equal, which indicates good reversibility of underlying electrochemical processes (Figure 13a). The maximum currents observed on the anode and cathode curves of the steel support are 1.5 μA (3.5 V) and 2.0 μA (3.0 V), respectively. According to [73], more detailed CV peaks can be observed using different scan rates.   Considering that lithium, silicon, iron, and nickel are present in the bulk of coating, many electrochemically active phases can be found and formed in it. The most probable phases include  [76]. The positions of the current peaks observed in the CV curves suggest that the charging and discharging capacities are accompanied by a change in the charge of nickel ions. Because NiO exhibits poor electrochemical activity in the range 2.5-4.3 V (as seen by the low increment of current in the CV curves shown in Figure 13a) it does not significantly contribute to the capacity of the coating. The impurities observed in the films and fewer peaks in the CV curves indicate that composition of films differs from that of LiNiO 2 .
To mitigate the effect of film thickness on the results of the study, the absolute values of the discharge capacities were divided by the volume (µm × cm 2 ). The LNO-1/10 samples annealed at 400-600 • C did not exhibit any discharge capacity. A moderate capacity was observed after thermal treatment at 700 • C. A quasi-linear decrease in voltage during discharge was observed (Figure 13b) for the LNO-1/10 700 • C sample, which is similar to the form of the capacitor discharge curve caused by the ion desorption (discharge) process. However, the shape of the CV curve ( Figure 13a) is closer to that of a cathode material than to that of capacitor electrodes. Therefore, a quasi-linear decrease in voltage may indicate that discharge occurs at relatively high currents.
The increase in the annealing temperature to 800 • C is accompanied by a change of discharge curve and capacity increase. In the discharge curve (Figure 13b), two sections can be distinguished, which differ in the angle of inclination. In the first section, a gradual decrease in voltage is observed during the discharge process; in the second, a sharp decrease in voltage occurs. This form of discharge curve is also observed during the discharge of powder cathode materials LiNi a Co b Mn c O 2 (where a + b + c = 1) and LiNiO 2 [77]. The subsequent increase in temperature to 900 • C did not result in a noticeable growth in discharge capacity. The electrochemically active phase is likely to be formed at 700-800 • C.
The presence of nickel atoms in the LNO-1/10 sample is confirmed by XPS and EDX analyses. Annealing stimulates the formation of the electrochemically active phase in the LNO-1/10 and LNO-M samples. According to the XRD results, only one nickel-containing phase (NiO) is manifested in the pattern of the LNO-M samples. Nevertheless, inclusions of the LiNiO 2 phase were observed by SAED. Considering the abovementioned results, it may be assumed that after calcination of LNO-1/10 and LNO-M, similar amounts of the electrochemically active phase are formed, which results in approximately equal discharge capacities. For all the studied samples, with an increment in the discharge current, a capacity decrease was observed (Figure 13c). At maximum discharge current (80 µA) the discharge capacity varied in the range 10-17 µAh·µm −1 ·cm −2 (41C). A subsequent decrease in discharge current led to an increase in discharge capacity, although the observed values were smaller than at the beginning of testing. Thus, the obtained ALD films can be discharged by relatively high currents, but the capacity decreases during cycling.
The capacity and specific capacity values of the model cathode material films were calculated assuming the following film parameters: thickness 1 µm, area 1 cm 2 , and true density of the cathode material, as well as the electrical test parameters: voltage range 4.2-2.8 V and nominal current discharge ( Table 5). The specific capacity of the films obtained in the present work varied in the range of 20-26 µAh·µm −1 ·cm −2 and lies in the interval of capacities observed for other ALD-deposited cathode materials, but approximately one-fourth of the expected value for LiNiO 2 (103 µAh·µAm −1 ·cm −2 ). For most ALD-deposited films of cathode materials, the specific discharge capacity is less than typical for model cathode material films of the same composition. The smaller capacities of deposited films may be due to the lower density, impurities, imperfect structure, and relatively high discharge currents.  [79] 10.58 54 -/0.1 C LiMn 2 O 4 [80] 10.07 51.4 -/1.0 C LiNiO 2 [77] 20.1 103 -/0.5 C

Conclusions
Atomic layer deposition of lithium-nickel-silicon oxide thin films was performed using LiHMDS and NiCp 2 as precursors and remote oxygen plasma as a counter-reagent. XRD, EDX, and XPS data indicated that LNO nanofilms deposited using the supercycle approach either do not contain Ni (LNO-1/1 and LNO-1/2) or only contain Ni in small amounts (LNO-1/3, LNO-1/10). However, these films contain lithium carbonates on the surface and a large amount of silicon in the bulk. The latter forms crystalline Li 2 SiO 3 and Li 2 Si 2 O 5 upon annealing at 800 • C. Obviously, silicon arises from LiHMDS, a silicon-containing precursor.
A multilayered LNO-M sample was successfully deposited using ALD of NO, LNO, LO, LNO, and NO layers. XRD data showed that the LNO-M sample contains crystalline NiO and annealing at 800 • C leads to the formation of Li 2 Si 2 O 5 and probably Li 5 Si 22 intermetallide. Local analysis of LNO-M-800 by SAED also showed the presence of LiNiO 2 . According to XPS and TOF-SIMS depth profiling, the annealing caused interdiffusion of layers, which leads to homogenization of the layer composition. In addition to homogenization, iron, chromium, and nickel diffuse from the stainless-steel substrate into the film. However, STEM analysis showed that the annealed films are not homogeneous at the micro/nanoscale. Submicron-scale crystallites of NiO/LiNiO 2 and predominantly amorphous silicon enriched layers were found.
Based on the shapes of the CV curves and the discharge curves, it can be assumed that the discharge capacity of deposited films is due to the intercalation of lithium ions, during which the charge of nickel ions changes. We assume that the LiNiO 2 detected by SAED is this electrochemically active phase. The values of the specific capacities for annealed LNO-M samples were in the range of 20-26 µAh·µm −1 ·cm −2 (at discharge currents of 5-7 C) and they are lower than that calculated for a dense LiNiO 2 film (0.5 C, 103 µAh·µAm −1 ·cm −2 ), but very close to 27 µAh·µAm −1 ·cm −2 , as obtained for ALD-deposited LiCoO 2 thin film.
Thus, using ALD method, the positive electrodes were prepared that can be discharged by relatively high currents and potentially suitable for creating thin-film lithium-ion batteries with increased power density. Considering that the ALD method allows coating on substrates with high aspect ratio trenches (high specific surface area), the specific weight of active material applied per unit of geometric area of the substrate can be augmented. As a result, the working time interval can be extended due to higher energy density.