Substrate Dependent Growth Behaviors of Plasma-Enhanced Atomic

Nov 28, 2012 - the phase of the films.2−4 Metal oxide (MO) films, such as NiO as in this ..... (at ncy of 0), respectively, suggesting that the larg...
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Substrate Dependent Growth Behaviors of Plasma-Enhanced Atomic Layer Deposited Nickel Oxide Films for Resistive Switching Application Seul Ji Song,† Sang Woon Lee,† Gun Hwan Kim,† Jun Yeong Seok,† Kyung Jean Yoon,† Jung Ho Yoon,† Cheol Seong Hwang,*,† Julien Gatineau,‡ and Changhee Ko‡ †

Department of Materials Science and Engineering and Inter-university Semiconductor Research Center, Seoul National University, Seoul 151-744, Korea ‡ Air Liquide, 28, Wadai, Tsukuba-Shi, Ibaraki Pref., 300-4247, Japan ABSTRACT: In this study, NiO thin films were deposited via a plasma-enhanced atomic layer deposition (PEALD) on metal (Pt, Ru, and W) substrates using a bismethylcyclopentadienyl-nickel ([MeCp]2Ni) precursor followed by a reaction with plasma-enhanced oxygen gas. The ALD temperature regime of NiO films was defined between 150 and 250 °C, while substrate temperature higher than this region induced the thermal cracking of precursors. The saturated PEALD rates of NiO film on Pt, Ru, and W substrates were 0.48, 0.58, and 0.84 Å/cycle, respectively, even though it has been usually regarded that the substrate effect on the saturated ALD rate vanishes after covering the entire surface with the growing films. At the initial stage of film growth, the NiO film showed enhanced nucleation behavior on the W and Ru substrates, whereas it did not show enhanced growth behavior on the Pt substrate. X-ray photoelectron spectroscopy revealed that the surface of a NiO film, which is thick enough for the W substrate not to influence the analysis, contains WO3 bonding states while the films grown on other metal substrates did not show any oxidation states of the substrate metal species. This could be due to the fact that the diatomic bond strength of W−O is stronger than that of Ni−O, which may induce the layer inversion during the ALD of NiO on the W substrate, and the surface W−O promotes the surface chemical reaction. This can result in the eventual increase of the saturated growth rate even in the ALD mode. The supply of oxygen to the adsorbing Ni-precursor by the reduction of a previously oxidized Ru substrate enhanced the initial growth rate of NiO film but this does not affect the steady-state growth rate on the Ru substrate. The small lattice mismatch between the NiO and Pt, as well as the identical crystal structure of the two materials results in the local epitaxial growth of NiO film on Pt substrate even though the growth temperature was only 250 °C. The NiO films on the W substrate showed reliable bipolar resistance switching in a wide temperature range (25−100 °C), which provides new opportunities for the next generation nonvolatile memory applications. KEYWORDS: NiO, W, Pt, Ru, O2 plasma, plasma enhanced atomic layer deposition, resistive switching memory substrate into the ALD of MO films. These films were grown on Ru metal, which is vulnerable to oxidation during the strong oxygen source (such as O3 or plasma-enhanced O2) pulse step and is also susceptible to the reduction during the precursor pulse step. In this case, the growth rate during initial tens of cycles was highly increased compared with the saturated steadystate growth rate, which is usually estimated from the slope of best linear fitting graph of film thickness (or mass gain) vs number of ALD cycles (ncy) when the film becomes thick enough not to be influenced by the possible reaction with the substrate. This was attributed to the chemical vapor deposition (CVD)-like growth of the films during the initial cycles due to the (unwanted) supply of oxygen to the precursor molecules during the precursor pulse step, where the oxygen was supplied by the reduction of the previously oxidized Ru substrate.5−7 This is basically due to the lower stability of the RuOx layer

I. INTRODUCTION The chemistry-specific (surface structure, precursor molecules, and reaction agent) saturated deposition rate (deposited amount per cycle) with respect to the precursor and reactionagent pulses is the key feature that characterize the atomic layer deposition (ALD).1 However, the surface chemistry in ALD is generally very complicated involving the surface reaction sites, gas phase molecules, reaction byproducts, and the exchange of electric charges between the surface and precursor molecules. An even more complicating factor could be the involvement of a substrate material or previously grown film material. Diffusion of a substrate material can influence the growth rate as well as the phase of the films.2−4 Metal oxide (MO) films, such as NiO as in this study, are also typical materials of which ALD rate are seriously influenced by these factors, especially when the films were grown on metal substrates of which oxides are unstable or stable against the chemical interaction with the incoming precursor molecules or growing films. Substrate enhanced growth of stable oxides, such as TiO25,6 and SrO,7 is a representative example of the involvement of a © 2012 American Chemical Society

Received: July 11, 2012 Revised: November 14, 2012 Published: November 28, 2012 4675

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of NiO film in detail for manufacturing nanoscaled conformal film as well as to improve the RS properties in ALD oxide films for NVM application. The electrode is another crucial factor that needs to be considered for the ALD NiO film to be used as a fluent RS material. For various reasons, RS is generally highly dependent on the types of electrode.17,18 Therefore, the study on the ALD behavior regarding NiO films on different electrodes and electrical properties of the NiO films is a crucial task. Ni is one of the difficult metals that could have a facile ALD precursor due to its relative large atom size and heavy weight. Therefore, stable precursors having stable and simple ligand structure, such as alkyl- or alkoxy-ligands, are rare. Perhaps the most feasible Ni-precursor is cyclopentadienyl (Cp-) based ones thanks to their volatility and high reactivity. ALD of NiO films using Cp-based precursors were reported by several groups.19−21 However, their research has been limited to examine the influence of deposition conditions, such as substrate temperature (Ts) and types of oxygen source on the crystal structure and composition of prepared films. Furthermore, detailed reports on the influence of metal electrodes and the RS property of the films are not made yet. Accordingly, in the present study, the authors adopted a PEALD process, giving a high growth rate and sufficient step coverage, in order to examine the ALD behavior of NiO films on different metal substrates and their electrical properties as the RS material.

compared with the growing oxide films. However, once the metal surface is completely covered with the growing film, and the growing films become thick enough, no more such oxygeninduced CVD-like growth is attained and the saturation growth rate is determined by the factors more inherent to ALD, such as molecule size, density of the surface adsorption site, and the reaction kinetics with oxygen, etc. On the other hand, if the substrate material is nonreacting, such as Pt, as shown in this work, or forms highly stable oxides, such as Si, such substrateenhanced growth is hardly observed, and even nucleation retardation is sometimes observed, which is generally an unwanted aspect of ALD. However, there have been cases seldom reported in ALD where the saturated steady-state growth rate was still influenced by the substrate material even when the growing films completely covered the metal surface and the films were thick enough to suppress the oxygen exchange through the films. In this work, the authors studied the substrate effects on the growth rate of NiO film via a plasma-enhanced ALD (PEALD) method, where the substrates are inert to the chemical reaction (Pt), interact with the growing film only at the interface (Ru) or keep interacting with the growing film throughout the whole growth cycles (W). Since the NiO films were grown simultaneously at a time by the PEALD on the three metal films, possible process variations can hardly influence the results whereby this could be a feasible model system to examine the complicated ALD mechanism regarding MO film depending on the types of metal electrodes. In addition, NiO film has been researched extensively as a resistive switching (RS) material for next generation nonvolatile memory (NVM) devices among the various binary transition metal oxides.8−10 As reported recently, the RS phenomenon in several MO films was triggered from the evolution of oxygen vacancy via the redox mechanism through the application of an electric field.11 For the specific case of NiO, when the conducting segments create a connected network, the electrical current drastically jumped to a high level, which is usually controlled to a compliance level in order to protect the sample from complete breakdown. Therefore, this is called electroforming, which actually coincides with a controlled soft breakdown. In contrast, the current was dropped down when the conducting network was disassembled by mostly the Joule heat effect. This RS phenomenon was closely correlated with a local nonstoichiometry within the oxide films.10,11 In many studies, NiO films were fabricated by reactive-sputtering or thermal oxidation of Ni metal. However, it was widely reported that the electrical properties of NiO film are highly sensitive to the deposition parameters such as oxygen partial pressure, plasma power, and deposition temperature, which generally make the stable production of fluently switching NiO films difficult.12−14 In this regard, the ALD method might be a viable solution to the reproducibility problem regarding the reactive sputtering process for NiO film thanks to its inherent self-limiting and saturating deposition behavior. In addition, it has an advantage in adjusting the film thickness and composition of films with atomic level accuracy by controlling the number of deposition cycles and varying the types of oxygen source or precursor/ reactant injection sequences. Moreover, it has unprecedented step coverage over severely topological surfaces, which is a crucial factor for the development of a vertically integrated device structure as an appropriate structure for a terabit-scale NVM.15,16 Therefore, it is necessary to study the ALD process

II. EXPERIMENT An 8-in.-scale single wafer shower-head type PEALD reactor (CN-1 Co., Plus-200) was used to deposit NiO films at wafer temperatures ranging from 150 to 300 °C. The bubbler for the Ni-precursors was heated to 50 °C and the precursor was delivered to the reaction chamber with Ar carrier gas at a flow rate of 50 sccm. Bis-methylcyclopentadienylnickel(II) [Ni(MeCp)2], synthesized by the Air Liquide company, having a vapor pressure of ∼0.1 Torr at the bubbling temperature was used as the Ni-precursor. The chamber pressure during ALD was ∼1.2 Torr. O2 gas was supplied with Ar gas during an oxygen source pulse step; each flow rate was 50 and 1000 sccm, respectively. The typical ALD process sequence consisted of a precursor pulse (3 s), Ar purge (5 s), oxygen source pulse (total 4 s while plasma power (100 W) was turned on during last 3 s), and Ar purge (5 s). In order to stabilize the partial pressure of O2 in the chamber, an O2 gas pulse step (1 s) was inserted before the plasma power was turned on (3 s). Various metal layers, such as sputter deposited-W (100 nm)/TiN (10 nm)/SiO2 (100 nm)/Si, sputter deposited-Pt (100 nm)/SiO2(100 nm)/Si, sputter deposited-Ru (30 nm)/Ta2O5 (10 nm)/SiO2 (100 nm)/Si were used as substrates. For comparison, 100-nm-thick electron beam evaporated-Pt film was also prepared. Here, the sputtered- and electron beam evaporated-Pt films are called sputter Pt and e-beam Pt. The physical thickness and the layer density of the grown NiO films were measured by an ellipsometer (Gaertner Scientific Co., L116D) and energy dispersive X-ray fluorescent spectroscopy (XRF, Themoscientific, ARL Quant’X). Film thickness and density were also checked by an X-ray reflectivity (XRR, PANalytical, X′pert Pro) measurement. The film surface morphology was examined by atomic force microscopy (AFM, JEOL, JSPM-5200) as well as scanning electron microscopy (SEM, Hitachi, S-4800). The chemical binding states of the films were determined by X-ray photoelectron spectroscopy (XPS, ThermoVG, SIGMAPROBE), and depth profiling was performed with Auger electron spectroscopy (AES, Perkin-Elmer model 660). The crystal structure of the NiO films was investigated by the X-ray diffraction (XRD, PANalytical, X′pert Pro) with the glancing angle scan mode (GAXRD) and the θ−2θ scan mode using Cu Kα Xray radiation. The electrical properties were measured using a semiconductor parameter analyzer (HP 4145B) and a pulse generator (HP81110A). For the electrical test, Pt top electrodes with an area of 4676

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55 000 μm2 were deposited by an electron beam evaporation method through a metal shadow mask.

activation energy on different substrates suggests that the detailed ALD reaction kinetics could be different, which was indeed the case, as shown later. Another notable finding is that the GPC value of NiO film is highly dependent on the types of substrate even in the ALD temperature regime. The films deposited on the W substrate showed much higher GPC values than those on the Ru and Pt substrates at all deposition temperatures. GPC of the NiO films on the Ru substrate is slightly higher than that of the Pt substrate at Ts < 250 °C. However, the difference becomes larger at a higher Ts. The possible reasons for such a substrate-dependent GPC are discussed in detail below. Self-limiting ALD saturation behaviors of the GPC with respect to the increase in the Ni-precursor pulse and purge times are confirmed at a Ts of 250 °C, as shown in Figure 2a and b. In Figure 2a and b, the precursor purge and pulse time was fixed at 5 and 3 s, respectively. The GPC was estimated by dividing the layer density by an ncy of 100, in this case. When the precursor pulse and purge time was >3 and >5 s, respectively, the GPC of Ni on Pt and Ru substrates was saturated, suggesting that the ALD-specific growth is well attained under this condition. The saturation behaviors of GPC with respect to the oxygen source pulse (plasma application time) and purge time were also confirmed (data not shown here). Nevertheless, the GPC of NiO film deposited on the W substrate is always higher than those of other substrates, and the complete saturation of GPC with respect to the precursor pulse time was not obtained even at 7 s. The reason for this can be understood from Figure 2b, where the purge time of 5 s is not sufficient under this condition; a purge time longer than 7 s is necessary to achieve a complete saturation. However, the difference in the GPC between a purge time of 5 and 7 s is small, so the following experiments were performed using the same purge time of 5 s for consistency. Figure 2c shows the typical XRR spectra of the NiO films grown on the different substrates, the thickness of which was estimated to be 36.9, 32.6, and 36.4 nm, respectively, on the Pt, Ru, and W substrates, which is quite close to the thickness estimated by ellipsometry (36, 34, and 37 nm). The simultaneously estimated film density was also indicated in the figure, where the NiO film on Pt showed the highest density, which may have a relationship regarding the distinct crystallographic structure, as shown later. The film on the W substrate showed the lowest density, which may be correlated with the highest growth rate mediated by the WO3 formation on the growing surface, as discussed in the following. The variations of the GPC of NiO films depending on the substrate types were further examined by estimating the variations in film thickness and layer density as a function of

III. RESULTS AND DISCUSSIONS III-1. Growth Behavior of the NiO Films Depending on the Substrates. Figure 1a and b shows the changes in the

Figure 1. Changes in the growth per cycle (GPC) of NiO films grown on a W, Pt, and Ru substrates, depending on the substrate temperature plotted in (a) linear scale and (b) Arrhenius form.

growth per cycle (GPC) of the NiO films grown on W, Pt, and Ru substrates, as a function of the Ts (150−350 °C) plotted in a linear and Arrhenius scale, respectively. Here, ncy was 300 on each substrate, and the GPC was calculated by dividing the estimated Ni layer density of each sample by 300. The GPC of the NiO films slightly increased up to ∼250 °C of Ts and then rapidly increased at higher Ts on W and Ru substrates, while it showed an increase much less evident in the higher Ts region on the Pt substrate. The transitions on each substrate were seen more evidently in Figure 1b, and such a transition usually ascribed to the change from the ALD-type deposition at lower Ts to the CVD-type deposition at the higher Ts region. The CVD-type deposition is believed to involve the thermal cracking of the precursor, which is not desirable for the wellcontrolled ALD type film growth. The relatively smaller activation energy for the CVD-like region on the Pt substrate (10.1 kJ/mol), compared with other cases (13.3−16.2 kJ/mol), could be ascribed to the highly oriented crystallographic structure of the NiO film grown on this substrate (highly 111direction preferred growth along the normal direction to the substrate surface), as shown later. ALD regime of NiO films could be defined below 250 °C, and the Arrhenius plot indicates more clearly that the transition from the process of ALD to CVD begins at 250 °C, as discussed in Figure 1b. Even in the ALD window, the thermally activated growth behavior is ascribed to the activation of the ligand exchange reaction with an apparent activation energy of 5.80−6.59 kJ/mol (Figure 1b), which was estimated from the slopes of the best-linear fitted graphs. The slightly different

Figure 2. Changes in the GPC of NiO film as a function of (a) Ni-precursor feeding time and (b) Ar purge time after Ni-precursor feeding step at the substrate temperature of 250 °C. (c) X-ray reflectometry measurement of NiO films on W, Pt, and Ru substrate. 4677

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ncy, as shown in Figure 3a and b. This is a feasible method to check if there are any incubation cycles or substrate-enhanced

suggesting that the GPC already reached the steady state growth at an ncy of ∼25−30 on these substrates. However, as shown above, the W substrate requires a much larger ncy of ∼100 in order to reach the steady-state GPC. In ALD, it is not unusual to encounter the strong influence of substrates on the film nucleation and growth depending on the type of substrates. However, it is generally believed that the substrate effect of the film growth vanished after covering the entire surface with the growing film, which is not the case in this study. Such peculiar growth behaviors of NiO films depending on the types of metal substrate layers are further examined, considering the possible diffusion of the substrate element into the growing film. First, the details of NiO ALD on the Pt substrate, which is believed to have minimal influence on the ALD considering its inert chemical nature, were examined. Figure 4a and b shows

Figure 3. Changes in (a) thickness of NiO film and (b) layer density of Ni as the number of cycles of ALD increased. Both layer density and thickness increased linearly with the ALD cycle number and the saturated growth rates of NiO films on W, sputtered-Pt and Si substrate were 0.84 (87.5), 0.58 (57.7), and 0.48 (49.1) Å/cycle (ng/ cm2/cycle), respectively. (c) Changes in film layer density on the various substrates (W, Pt, and Ru) below 50 cycles and (d) the derivative of the layer density in (c) as a function of ALD cycle number.

growth of the films. Here, the films were deposited at a Ts of 250 °C and the Ni-precursor pulse/purge and oxygen source pulse/purge times were 3/5 and 3/5 s, respectively, with ncy ranging from 50 to 800. Figure 3c shows the variations in the layer density of the films on the three substrates as a function of ncy where it ranges from 1 to 50, which shows the influence of substrate on the NiO growth more clearly. It can be understood that the NiO films grew fluently without any nucleation retardations on the Ru and W substrates but showed a slight retardation of the GPC up to ncy of ∼3 on the Pt substrate. The NiO on the W substrate shows a largely enhanced GPC up to an ncy of ∼3, and it decreases abruptly after an ncy of 5 (Figure 3d), suggesting that the W substrate heavily enhanced the nucleation of NiO film. Between ncy of 5 and 50, the GPC on the W substrate in terms of the layer density/cycle was approximately 110−120 ng/cm2/cycle (∼1.2−1.3 Å/cycle), which is higher than that obtained from the slopes of the best linear fitted graph shown in Figure 3b (87.5 ng/cm2/cycle, and 0.84 Å/cycle in Figure 2a) obtained in the ncy range >100. This clearly shows that the W substrate enhances the growth of NiO film. The slightly retarded GPC on the Pt substrate during the very first few cycles increases to an ncy value of ∼20−25 and becomes saturated afterward. Such a variation in the GPC on Pt is consistent with the nucleation-induced roughening and coalescence model suggested by R. L. Puurunen.22 The Ru substrate also shows a slight substrate-enhanced GPC up to an ncy of 25−30, but then, the GPC decreases to the steady-state value afterward in Figure 3d. The steady-state GPC shown in Figure 3d for the Pt and Ru substrates after an ncy of ∼30 is similar to the values obtained from the best linear fitted graphs shown in Figure 3b (49.1 ng/cm2/cycle and 0.48 Å/cycle for Pt; 57.7 ng/cm2/cycle and 0.58 Å/cycle for Ru) for ncy > 100,

Figure 4. Plane view SEM images of NiO films grown on (a) sputtered-Pt and (b) e-beam evaporated-Pt substrate. (c) Change in 10-point median roughness (Rz) obtained by AFM of NiO films on the W, sputtered-Pt, and e-beam evaporated-Pt substrate with increasing ncy. (d) Changes in layer density and thickness of NiO films deposited on sputtered-Pt and e-beam evaporated-Pt substrate as a function of number of cycles. XRD patterns in (e) glancing angle and (f) θ−2θ scan mode of 30 nm-thick NiO films deposited on W, Pt, and Ru substrate, respectively. (The diffraction peaks of NiO marked by black circle.)

the plain view SEM images of the 40-nm-thick NiO film grown on the sputter Pt and e-beam Pt substrates, respectively. The film shows quite a rough morphology, which resembles the original morphology of the Pt substrate itself (SEM images are not shown here). This means that the NiO film was grown along the grain of the Pt electrode. This can be closely correlated with the specific crystallographic growth behavior shown in Figure 4e and f. Since such a rough surface morphology of the substrate can influence the ALD behavior, the possible correlation between the roughness and layer density of NiO film was examined using the two types of Pt substrate: sputtered-Pt and e-beam Pt. Figure 4c shows the 4678

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Figure 5. Cross section and plane view SEM images of NiO films grown on (a, b) W and (c, d) Ru substrate. The ALD cycles repeated by 700.

was found in this diffraction spectrum near 2θ ∼ 37°, which is even stronger than that of the Ru substrate peak shown in the same graph. This strongly implies that the NiO film grows locally epitaxially due to the favorable lattice constant match (lattice mismatch was ∼6% between 0.4177 and 0.3923 nm for NiO and Pt) with the same face centered cubic structures regarding these two materials. The crystal structures of W and Ru are body-centered cubic and hexagonal structures, respectively. Due to this growth behavior, the NiO layer grew very densely, as can be understood from the highest density revealed by the XRR results shown in Figure 2c. The lowest steady-state saturated GPC of NiO film on Pt compared with other cases appears to be correlated with this specific growth behavior. The local epitaxial growth may restrict the structure and distribution of surface chemisorption sites, which can reduce the GPC. GPC of a certain ALD process is generally governed by the surface density of the available chemisorption sites and steric hindrance effect originating from the size of the precursor molecules.23−25 For the given precursor molecules, and the substrate surface being covered with the growing films, the density of chemisorption sites may have a larger control regarding the saturated GPC. It is believed that the density of surface chemisorption sites depends on the crystallographic surface structure of the growing NiO film, the structure of which is rock salt. It must also be noted that the crystallographic surface exposed to the gas phase is not (111) planes for the highly 111-preferred NiO film on the Pt substrate as can be understood from the triangular shape of the Pt grains with tilted side planes, which are usually comprised of (100) and (110) type planes.26 Therefore, it is believed that the adsorbing Ni-precursor molecules find their chemisorption sites mostly on the (100) and (110) type planes of NiO, of which adsorption site density is lower than that of the (111) type plane. It should also be noted that this phenomenon is different from the situation of conventional epitaxial growth of films where the precursor molecules are already thermally decomposed and the adsorbed surface atoms migrate until they find stable positions, such as ledge or kink sites.27 Here, due to the partly retained ligands of the adsorbing molecules, which exert the steric hindrance effect, the available adsorption site density mainly governs the GPC. On the other hand, the NiO films grown on other substrates grew with no specific preferential direction, meaning that the (111) type plane can contribute to the precursor adsorption. This may correspond to the higher saturation of GPC on the Ru substrate compared to the Pt case.

changes of a 10-point median roughness (Rz) of NiO films grown on different substrates, as a function of ncy, determined by measuring the five highest peaks and lowest valleys over the sampling area of 500 × 500 nm2 using the AFM. Here, the data on W substrate was also included for comparison. The Rz/ average grain sizes of sputter Pt, e-beam Pt, and W substrates are 13.4/87.4 ± 8.5, 8.39/29.4 ± 4.7, and 11.8/36.2 ± 7.8 nm (at ncy of 0), respectively, suggesting that the larger grain size resulted in the higher roughness. It can be understood that the Rz value of NiO film on the sputter Pt substrate does not change largely but that of the film on e-beam Pt rather increases with the increasing ncy and becomes similar to that of the film on the sputter Pt case at ncy of 300. This confirms that the NiO film was grown along the Pt grain morphology with possible local epitaxy. On the other hand, the Rz value on W substrates decreases abruptly with an ncy of up to ∼50 and remained at a low value after such an ncy suggesting that the NiO nucleates preferentially at the valleys and coalesces on this substrate. Figure 4d shows the variations of layer density and thickness of the NiO films grown on of the sputter Pt and e-beam Pt substrates as a function of ncy. Although the e-beam Pt substrate had a lower roughness, so a lower surface area, the NiO films generally have a slightly higher layer density and thickness suggesting that the surface topology of a substrate is not the major factor causing the difference in saturated growth rate of NiO film. This also suggests that the much higher GPC on the W surface has little relevance with the surface topography, even though W has a rougher surface morphology. The GAXRD and θ−2θ scan XRD data shown in Figure 4e and f, respectively, show another characteristic feature of the NiO film grown on the Pt substrate. It must be noted that the GAXRD cannot detect the diffraction peak from the crystallographic planes, which are parallel to the substrate surface, while θ−2θ scan XRD detects only crystallographic planes, which are parallel to the substrate surface due to their specific relative geometry between the detector and reciprocal lattice during the measurements. The NiO film on the sputter Pt substrate does not show any peak, while those on the Ru and W substrate show clear diffraction peaks corresponding to the cubic NiO phase in the GAXRD (Figure 4e), suggesting that the NiO film on the Pt electrode is either amorphous or crystallized with a highly preferred orientation. However, Figure 4f shows a very strong Pt (111) peak near 2θ ∼ 40°, suggesting that the Pt film has a very strong textured structure with (111)-preferred growth behavior. Interestingly, a very strong NiO (111) peak 4679

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The surface atomic densities of (100), (110), and (111) plane of face-centered cubic structure are 11.46, 8.106, and 13.24 atoms/nm2. A higher surface atomic density coincides with a higher chemical adsorption site density. Figure 5 show the cross-section and plain view SEM images of the 65-nm and 43nm-thick NiO films on the W and Ru substrates, respectively, showing the very distinctive morphology of the film from that on the Pt substrate. The fine microstructure coincides with the smoother surface morphology revealed by the AFM. However, the far higher GPC of the NiO film on the W substrate compared with the other two cases cannot be explained by this effect. On the Ru and W substrates, the GPC of the NiO films appears to be divided into two and three regions (Figures 3c and d), respectively. On the Ru substrate, the GPC was higher for an ncy up to ∼25 and reaches to the steady-state value afterward. On the W substrate, however, it shows a very high value up to ∼3 cycles and an intermittent high value up to ∼100 cycles and then reaches the steady-state value afterward. It must be reasonable to assume that the higher GPC during the initial cycles are closely correlated with the chemical interaction between the substrate and Ni-precursor. The increased GPC during the initial ALD cycles on the Ru substrate has been reported for other oxides, such as TiO25 and SrO,7 when the adopted oxygen source had a strong enough oxidation potential to oxidize Ru. In those cases, the oxygen sources oxidized the Ru layer to RuO2 or RuOx (x < 2) during the oxygen source pulse step, which then reduced to Ru during the subsequent metal-precursor pulse step due to the weaker binding energy between the Ru and O compared with that between Ti and O or Sr and O. This provides the incoming metal-precursors with highly reactive oxygen atoms, which can induce even multilayer oxide formation during the metalprecursor pulse step although the ALD was meant to induce (sub)monolayer formation of the partly decomposed metalprecursor adsorption layer at this stage. This results in the much higher GPC as long as this oxygen supply through the very thin oxide layer deposited is maintained (oxygen must diffuse into the underlying Ru during the oxygen source pulse step and diffuse out to the growing film surface during the metal-precursor pulse step). Once the growing film becomes thick enough to suppress these oxygen diffusions, the substrate enhancement effect disappears and a normal steady-state growth rate is achieved. The same mechanism can be invoked here in order to explain the initial enhancement of GPC on the Ru substrate. Choi et al. reported that plasma-activated O2 gas has a strong enough potential to oxidize Ru electrode, as O3 does, which cannot be expected from H2O.28 Therefore, it is also probable that the Ru electrode is oxidized during the plasma-activated O2 gas pulse step. Since Ni has a higher oxidation potential than Ru, the oxidized RuOx must be reduced to Ru providing the adsorbing Ni-precursors with oxygen atoms during the precursor pulse step, which eventually increases the GPC, as shown in Figures 1 and 3. As the NiO film becomes thick enough to suppress both the diffusion-in and diffusion-out process of the oxygen atoms, the GPC recovers the normal saturation value determined by the precursor chemistry and surface adsorption site density. Further confirmation for such a mechanism was provided in Figure 6a and b, where the GAXRD patterns of the RuO2 substrate before and after seven cycles of Ni-precursor injection without the plasma-activated oxygen gas pulse (Figure 6a) and

Figure 6. (a) GAXRD patterns of after seven cycles of Ni-precursor feeding on RuO2 substrate. (c) The initial growth behaviors of NiO grown on W and O3-treated W substrate. XPS spectra of (b) Ru 3d5/2 and (d) W 4f7/2 binding energy of before/after seven cycles of source feeding on the RuO2 and O3-treated W substrate, respectively. The photoemission intensities were normalized to the same height.

XPS data of the same samples (Figure 6b) are shown. Here, the 20-nm-thick RuO2 film was prepared by a CVD process using the Ru (EtCp)2 precursor and O2 gas at a Ts of 250 °C. The diffraction peaks from the RuO2 disappeared completely and the peaks from the Ru showed up in 2θ range 42−44° just after the seven cycles of Ni-precursor pulses. XPS showed consistent results; the Ru 3d5/2 binding energy (BE) shifted from 281.0 (RuO2) to 280.1 eV (Ru) after the same treatment of the RuO2 film. Therefore, the Ni-precursor reduced the RuO2 to Ru, and the NiO film was formed even without subsequent oxygen source supply. In fact, after just seven cycles of Ni-precursor pulses, the sample showed a significant amount of Ni in the XRF measurement even though the plasma-activated oxygen pulse was not applied. Again, this is due to the lower oxide stability of RuO2 (ΔGf; −217.86 kJ/mol at 500 K, which corresponds to −72.62 kJ/atom mol) compared with the NiO (−192.84 kJ/mol at 500 K, which corresponds to −96.42 kJ/ atom mol) and high reactivity of the precursor. Lee et al. had reported a similar growth behavior of SrTiO3 films on the IrO2 and RuO2 substrates.5 In contrast, the same mechanism can hardly explain the substrate-enhanced growth behavior of the films on the W substrate. Even at room temperature, the W substrate would be easily oxidized since the oxidation potentials of tungsten oxides (WO2 or WO3) are high enough (−496.4 kJ/mol (−154.8 kJ/ atom mol) and −711.2 kJ/mol (−177.8 kJ/atom mol) at 500 K), to form a native oxide at the substrate surface. Furthermore, it had been reported the combustion products, such as H2O and CO2, were observed during the overall reaction of NiO ALD using a Cp-type precursor, which may further induce the oxidation of W.21,29,30 Due to this higher oxidation potential of WO2 and WO3 compared with that of NiO, it is unlikely that the same mechanism, that is, oxidation of the W substrate during the oxygen source pulse step and the reduction of it releasing oxygen atoms to the adsorbing Ni-precursors during the subsequent Ni-precursor pulse step, works and explains the 4680

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abrupt. The more graded concentration profile at the interface for the case of the W substrate compared with the Ru is due to the higher roughness of the W substrate. The NiO film on W substrate generally has higher oxygen content compared with that on the Ru substrate, which must be related with the diffusion of W onto the growing surface as discussed below. It must be noted that the absolute atomic concentration shown in Figure 7a and b are not accurate enough to tell the Ni/O ratio due to the lack of the appropriate reference sample, so the relative ratio must be relevant here. An even more notable finding can be made in Figure 7c and d, which shows the W 4f7/2 and Ru 3d5/2 XPS spectra taken from the surface of the 30nm-thick NiO films grown on the W and Ru substrates, respectively. The XPS shows a clear signature of WO3 on the NiO film surface regarding the W substrate whereas that of the NiO film on the Ru substrate does not show any presence of (oxidized) Ru. The strong peak near the BE of 285 eV is ascribed to the adventitious carbon, which was used to calibrate the BE of the spectra. Since the NiO film thickness is large enough than the escape depth of photoelectron not to detect the W substrate in the XPS analysis, the WO3 signal originates evidently from the film surface. However, the AES depth profile showed that the bulk of the NiO film does not contain a significant amount of W. Therefore, it can be inferred that a small amount of W was diffused into the thin NiO film and these diffused W atoms oxidized and remained on the film surface during the entire ALD process. This is believed to be the origin for all abnormal GPC behavior of the NiO film on this substrate. The higher the GPC of NiO on W, where actual surface is more like WO3, is basically ascribed to the higher density of adsorption sites on the WO3 surface compared with NiO. The adsorption sites for the Ni-precursor must be basically O-sites of surface oxides, where the hydrogen or any other reaction group bound to the oxygen mediates the chemical adsorption via the ligand exchange reaction during the precursor pulse step, although the actual adsorption could also be influenced by the size effect of the incoming molecules (steric hindrance effect). However, the fairly high growth rate (∼1 Å/cycle) in the intermediate high GPC region suggests that the steric hindrance effect has a relatively minor effect compared with the effect of the density of surface adsorption sites. Since the growing NiO film crystallizes into the cubic structure, its surface oxygen ion density can be calculated for any major crystallographic surface. However, the surface structure of very thin WO3 (or even WOx, which may coincides with the native oxide) is not known, so that the quantitative comparison for the surface oxygen ion densities of WO3 and NiO is not possible. However, it is quite evident that the W atom has a higher density of surface oxygen compared with the Ni atom considering its far higher valence state (+6) compared with that of Ni (+2). Therefore, it must be a safe hypothesis that the WO3 surface can provide a higher density of sites than the NiO surface for the precursor adsorption. When the very first cycles of the ALD process were performed on the W substrate, the Niprecursor molecules absorb on the native WO3 (or WOx) most likely via the mediation through the surface OH groups, which must be abundantly present on the native tungsten oxide. This must result in the highest GPC considering the fully WO3 (or native oxide) surface, as shown in Figure 3d. When the surface is covered with NiO, the possible adsorption site density decreased due to the higher portion of surface NiO, which corresponds to the decreasing tendency of GPC up to an ncy of

GPC enhancement during the initial growth on the W substrate. This can be supported by the following experiments. Figure 6c shows the variation of layer density as a function of ncy up to 15 on the W and O3-oxidized W substrate. Here, the data for the W substrate was reproduced from the data shown in Figure 3c, and the O3-oxidized W substrate was prepared by exposing the surface of the W substrate to O3 ambient (O3 concentration; 300 gm−3, 30 s at 250 °C) immediately prior to the ALD of NiO film. Figure 6d shows the XPS spectra of this O3-oxidized W substrate (black line) showing that a thin WO3 layer was formed. The thickness of the WO3 layer was estimated to be very thin (100. When the NiO film was thin enough not to significantly block the W-diffusion, such as ncy < ∼100, which coincides with an approximate NiO film thickness of ∼10 nm, a large amount of W diffused from the W substrate onto the growing film surface, so that the surface WO3 density could be high resulting in the higher GPC in that ncy range. However, when the films became thicker, such as ncy > 100, the diffusion from the bulk W layer becomes unlikely, and only surface WO3 molecules, which were in direct contact with the incoming Niprecursor molecules may works as the GPC enhancement factor. As can be understood from the tiny change in the XPS peak shape in Figure 6d, the surface WO3 molecules appear to be slightly reduced by the chemical interaction with the 4682

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ratio of 1:5 (The opening diameter and depth of the hole were each 70 and 350 nm.) was ∼85% as shown in the cross-section SEM image shown in Figure 9b. These results proposed that the self-limited ALD reaction takes place all over the surface of a hole under this deposition condition. III-2. Electrical Properties of the NiO Films. The electrical properties of NiO film grown on W substrates are discussed below. These correspond to only preliminary characterization results in order to show that the prepared films are of useful quality as the RS material. More in-depth and thorough characterization results with memory cell compatible sample geometry will be reported elsewhere. It was found that the NiO films on the W substrate showed the most plausible electrical properties due to its lower leakage at the pristine state, as can be understood from the current− voltage (I−V) curves of the pristine 40-nm-thick NiO films on the Pt and W substrates shown in Figure 10a when the applied

precursor molecules during the precursor pulse step, which eventually forms a Ni−O−Cp layer on top. This backs up the hypothesis shown in Figure 8. Then, the W atoms belonging to these defective WOx molecules may want to move onto the Ni−O layer surface to be more completely oxidized during the subsequent oxygen source pulse step. This recovers the surface WO3 containing the NiO surface structure, where the WO3 density must be lower compared with the thinner NiO case, in which the W substrate can still provide a certain number of W atoms onto the surface. Now, the growth results for the O3-treated W substrate shown in Figure 6c are reconsidered. The O3-treated W substrate is devoid of abundant surface OH groups, so that the Ni-precursor can rarely find anchoring sites when they were first pulsed on that substrate. Perhaps only a small amount of Ni-precursors chemically adsorb on the surface via the mediation by the trace amount of OH or the small chemical interaction with the WO3. Once this has occurred, the next plasma-enhanced oxygen pulse may induce additional surface adsorption sites to form, such as formates or carbonates30,31 on the Ni−O molecules or even on the WO3 surface. Then, the next ALD reaction may take place via the mediation by these secondary reactive species, which may explain the low growth rate shown up to an ncy of 5 (Figure 6c). Meanwhile, the bulk of the WO3 layer may remain intact up to this ncy, so that the diffusion of W atoms from the underlying W substrate onto the surface is not likely and no GPC enhancement can be achieved under this circumstance. However, as discussed above, the repeated supply of the Ni-precursor induced the reduction of some portion of the WO3 and the W atoms from these defective WO3 started to move onto the surface, and finally, an enhanced GPC was achieved. This surface WO3 density might be corresponding to the case with thicker NiO film on the nontreated W substrate, so that the GPC enhancement might be smaller than the W case in the same ncy. However, this cannot be unequivocally confirmed with the limited number of deposition experiments shown in Figure 6c although the data showed a slightly smaller GPC compared to the other case between ncy of 7 and 15. Figure 9a shows the layer density of Ni and thickness of NiO films at various locations within the 4- and 8-in.-diameter Si

Figure 10. (a) Current−voltage curves of the pristine NiO films grown on the W (closed) and Pt (open) substrates. RF-plasma power was varied between 100 W (black square symbol) and 300 W (red circle symbol). (b) XPS spectra from the valence band edge of NiO films on W substrate deposited with an RF-plasma power of 100 and 300 W, respectively. Typical resistive switching current−voltage characteristics of NiO films deposited on (c) W (bipolar) and (d) Pt (unipolar) substrates.

RF power during the oxygen source pulse step was 100 W. Therefore, samples on W substrates are taken for their electrical characterization. However, even on the W substrate, the RF power plays a crucial role in determining leakage current as can also be seen from Figure 10a. When the RF power was increased to 300 W, the NiO film becomes even leakier than that on the Pt with the RF power of 100 W probably due to the damaging effect of the harsh plasma condition. Figure 10b shows the valence band spectra of NiO films in XPS. It can be understood that the film with the higher RF power has a higher tail state density near the valence band edge into the band gap, which may induce the trap-mediated current conduction. Therefore, the 40-nm-thick NiO film grown with an RF power of 100 W on the W substrate was taken to check if a stable RS property can be achieved from this PEALD processed NiO films. Figure 10c shows the switching I−V curves with a bipolar type RS (BRS) mode where the top Pt electrode was biased

Figure 9. (a) Thickness and layer density of Ni at various locations within 4 and 8-in. diameter Si wafer. (C, center; T, top; R, right; L, left; B, bottom) ALD sequence of NiO repeated with 250 cycle. (b) Ni-precursor feeding time and (d) source purge time at the substrate temperature of 250 °C. (b) Cross-section SEM image of NiO film deposited on a contact hole structure, giving an aspect ratio of 1:5.

wafers. ALD sequences were repeated for 250 cycles at the Ts of 250 °C, and the thickness nonuniformity ((max − min)/2 average) of NiO film showed 0.5% and 4.4% in 4- and 8-in.diameter wafers, respectively. Furthermore, even the ALD process of NiO film consisted of a plasma sequence, the step coverage of film inside a contact-hole structure with an aspect 4683

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Notes

and the bottom W electrode was grounded during the measurement at different temperatures (25−100 °C). The sample shows feasible BRS performance in such a wide temperature range although the resistance ratio between the high and low resistance state is rather small (∼10) especially at high measurement temperatures. Even though the NiO film on Pt electrode showed a higher leakage current than that on W electrode, the NiO film with symmetrical electrode configuration (Pt/NiO/Pt) showed the unipolar switching mode at room temperature, as shown in Figure 10d, where only positive bias polarity was used to induce both set and reset switching. In this case, the resistance ratio was higher, as reported by many other previous works.32,33 Although the preliminary data presented in this section are not sufficiently detailed, the NiO films grown by this PEALD process showed a feasible electrical performance as the RS material in future resistive memory devices.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This study was supported by the National Research Program for the Nano Semiconductor Apparatus Development sponsored by the Korea Ministry of Knowledge and Economy (10034831) and the Converging Research Center Program through the National Research Foundation of Korea (NRF) (2011K000610).

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IV. CONCLUSION In this work, the influences of metal substrates, Pt, Ru, and W, on the growth behavior of NiO thin films were examined when they were deposited via a plasma-enhanced atomic layer deposition using a bis-methylcyclopentadienyl-nickel ([MeCp]2Ni) precursor and plasma-enhanced oxygen gas. The ALD should be performed in the temperature region between 150 and 250 °C, while a substrate temperature higher than this region induced the thermal cracking of precursors. The saturated steady-state PEALD rates of NiO film on Pt, Ru, and W substrates were 0.48, 0.58, and 0.84 Å/cycle, respectively, while the Pt substrate slightly retarded the initial nucleation of the NiO film, but the Ru and W enhanced the nucleation and growth of the film during the initial tens of cycles. The small lattice mismatch between the NiO and Pt and the identical crystal structure of the two materials resulted in the local epitaxial growth of NiO film on the Pt substrate even though the growth temperature was only 250 °C. The surface of the growing NiO film comprised of planes other than (111), which is believed to have the highest adsorption site density, on the Pt substrate, so that the steady-state growth rate is lower than the Ru substrate case where no specific growth direction was found. The steady-state growth rate of NiO film on the Ru electrode is believed to be the value that was free from any influence of substrate. The faster growth rate of the film on Ru during the initial cycles is ascribed to the CVD-like growth by the supplied oxygen atoms from the reduced RuO2 layer, which was formed during the previous oxygen source pulse step. On the other hand, X-ray photoelectron spectroscopy revealed that the surface of a 30-nm-thick NiO film contains WO3 bonding states while the films on other metal substrates did not show the oxidation states of the substrate metal films. This could be due to the layer inversion during the ALD of NiO on the W substrate, which must originate from the stronger diatomic bond strength of W−O than that of Ni−O. The surface WO3 can induce a higher growth rate because it can provide the surface with the higher density of adsorption sites for the incoming Ni-precursors. The NiO films on the W substrate showed reliable bipolar resistance switching in a wide temperature range (25−100 °C) as well as fluent unipolar switching at room temperature.

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