Atomic Layer Deposition of Iridium Thin Films Using Sequential

Article pubs.acs.org/JPCC

Atomic Layer Deposition of Iridium Thin Films Using Sequential Oxygen and Hydrogen Pulses Miika Mattinen,† Jani Ham ̈ al̈ aï nen,† Marko Vehkamak̈ i,† Mikko J. Heikkila,̈ † Kenichiro Mizohata,‡ ‡ Pasi Jalkanen, Jyrki Raï san̈ en,‡ Mikko Ritala,*,† and Markku Leskela†̈ †

Laboratory of Inorganic Chemistry, Department of Chemistry, University of Helsinki, P.O. Box 55, FI-00014 Helsinki, Finland Division of Materials Physics, Department of Physics, University of Helsinki, P.O. Box 43, FI-00014 Helsinki, Finland

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ABSTRACT: Atomic layer deposition (ALD) is an advanced thin-film deposition method based on self-limiting surface reactions that allows for the controlled deposition of conformal, high-quality thin films of various materials. In this study, we aimed to explore how modifying the deposition chemistry affects the growth and properties of iridium films. We demonstrated a new ALD process using sequential pulses of iridium acetylacetonate [Ir(acac)3], oxygen (O2), and hydrogen (H2) and compared this to the established Ir(acac)3 + O2 process in the wide temperature range of 200−350 °C. A reaction scheme is proposed to explain how both oxygen and hydrogen affect the film growth. Comprehensive information on film properties was obtained for both processes. In particular, the strong (111) texture seen in this study has not been reported before for ALD iridium films. Changes in film properties, especially lowered resistivity, stronger (111) texture, and faster nucleation compared to the Ir(acac)3 + O2 process, should motivate further studies on O2 + H2 processes of platinum metals.

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INTRODUCTION Atomic layer deposition (ALD) is a gas-phase thin-film deposition technique that relies on sequential, self-limiting surface reactions. These are realized by pulsing precursors onto a substrate alternately, separated by evacuation or purge periods, to avoid reactions in the gas phase that occur in the closely related chemical vapor deposition (CVD) technique. Self-limiting reactions allow for the deposition of pure, highquality films with excellent conformality and precisely controlled film thickness and composition.1−3 Platinum metals (Ru, Rh, Pd, Os, Ir, and Pt) are elements with outstanding physical properties, such as high melting point, low coefficient of thermal expansion, and low resistivity. Chemically, they have low reactivity and are thus resistant against oxidation; yet, they are also known for their high activity in catalysis. Notably, these metals are very rare in Earth’s crust, which, combined with their excellent properties and wide range of applications, makes them expensive. 4 Atomic layer deposition is an ideal method for depositing nanoparticles and thin films of platinum metals because of its superior controllability, which would possibly allow for the use of smaller amounts of these expensive elements.5,6 Although ALD of elemental materials has been challenging in general, working processes for all of the platinum metals have been found.7 Most of these processes use molecular oxygen (O2), a reactant that is otherwise usually found to be inert under typical ALD conditions, to combust ligands of the metal precursors. Successful use of molecular oxygen relies on its © 2016 American Chemical Society

dissociative adsorption on catalytically active platinum-metal surfaces, transforming it into the much more reactive atomic form.8,9 Most of the ALD studies on platinum metals have concerned platinum and ruthenium (see ref 7 for a thorough review on platinum-metal ALD), whereas iridium has gained less attention despite its attractive properties such as high melting point, high density, low electrical resistivity, and excellent chemical resistance.4 Previously reported ALD processes of Ir most often use O2 as a reactant with Ir(acac)3,10,11 (EtCp)Ir(COD),12 or (MeCp)Ir(CHD)13 at temperatures above 200 °C, whereas consecutive O3 and H2 pulses can be used to deposit Ir at temperatures below 200 °C.14,15 Plasma-enhanced ALD processes using H2 plasma,16 NH3 plasma,17 or mixed O2−H2 plasma18 have also been demonstrated. ALD iridium films have been used in optics and catalysis, as well as diffusion barrier and seed layer applications, for example (see ref 13 and references therein). In the present study, a new three-step ALD process comprising Ir(acac)3, O2, and H2 pulses, hereafter called the O2 + H2 process, is demonstrated and characterized. The process was inspired by the low-temperature Ir(acac)3 + O3 + H2 process, where the hydrogen pulse removes oxygen from the surface, transforming the film from iridium oxide to metallic Ir.15,19 Although the H2 pulse is not necessary for obtaining Received: May 3, 2016 Revised: June 28, 2016 Published: June 29, 2016 15235

DOI: 10.1021/acs.jpcc.6b04461 J. Phys. Chem. C 2016, 120, 15235−15243

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The Journal of Physical Chemistry C

θ−2θ diffraction geometry and Cu Kα (λ = 1.54 Å) X-ray beam. Rocking-curve measurements for film texture studies were performed from the Ir(222) peak to avoid the sloping background of the Ir(111) peak. X-ray reflectivity (XRR) measurements for the estimation of film roughness were performed with the same instrument. Film resistivity was determined based on four-point probe measurements (CPS probe station connected to a Keithley 2400 SourceMeter) and thicknesses measured by EDX. The uncertainty of resistivity was estimated based on the propagation of uncertainty using 5% error for the thickness and the standard deviation from several measurement locations for resistance. Impurities were analyzed by time-of-flight elastic recoil detection analysis (TOF-ERDA) using a 79Br7+ ion beam with an energy of 32 MeV.26

metallic films when O2 is used as a reactant, removing the adsorbed oxygen before the next Ir(acac)3 pulse affects film growth and properties. Compared to the O3 + H2 process, the reaction mechanisms are expected to be different because of the large difference in oxidation power of O3 and O2, as well as the different temperatures used. For other platinum metals, similar O2 + H2 processes were previously briefly tested using (MeCp)PtMe320 and Ru(EtCp)221,22 precursors. However, we believe that this is the first study considering film properties obtained with an O2 + H2 process, as well as the first demonstration of such a process for iridium. We characterized films comprehensively with respect to their crystallinity, morphology, composition, and resistivity. Thus, in addition to the new O 2 + H 2 process, new information on the corresponding O2 process has also been obtained. We believe that the present results can also be valuable for increasing the understanding of platinum-metal ALD processes in general.

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RESULTS AND DISCUSSION We began by studying how adding a hydrogen pulse to the Ir(acac)3 + O2 process affects the film growth at 250 °C. The H2 pulse length was varied between 0 and 5 s, and saturation of the growth rate was observed when the pulse length was at least 1 s (Figure 1), leading to a ∼20% lower growth rate compared to the O2 process (represented by 0-s H2 pulse).

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EXPERIMENTAL SECTION Iridium films were deposited using a commercial, flow-type, hot-wall ALD reactor (F120, ASM Microchemistry). Nitrogen (N2, AGA, 99.999%) served as both the carrier and purge gas. Continuous N2 flow was set to 400 sccm, leading to a pressure of about 10 mbar inside the reactor. Films were deposited on 5 × 5 cm2 silicon (100) substrates, which were blown clean with pressurized nitrogen before being loaded into the reactor. Native silicon oxide was not removed. Iridium acetylacetonate [Ir(acac)3, ABCR, 99%] was sublimed from an open glass boat heated to 155 °C inside the ALD reactor. Oxygen (O2, AGA, 99.999%) and hydrogen (H2, AGA, 99.999%) flows of 10 sccm were set using a needle valve and a mass flow meter. Precursors were pulsed onto the substrates alternately using inert gas valving and solenoid valves, separated by N2 purges. The deposition temperature was varied between 200 and 350 °C. In most cases, the purge and pulse times allowing self-limiting growth were fixed at 1 s based on the present results and the previous study by Aaltonen et al.10 Scanning electron microscopy (SEM, Hitachi S-4800) and atomic force microscopy (AFM, Veeco Multimode V) were used to study film morphology. Tapping-mode AFM imaging was performed in air using silicon probes (Bruker) with a nominal tip radius of less than 10 nm. AFM images were flattened to remove artifacts caused by sample tilt and scanner nonlinearity. Film roughness values and their error bars represent average values and standard deviations, respectively, of root-mean-square roughness (Rq) determined from at least three images for each sample. Image analysis was performed using the Bruker Nanoscope 1.5 program. Lateral grain size was estimated from the AFM images using the Watershed algorithm implemented in the Gwyddion 2.36 program.23 Transmission electron microscopy (TEM, FEI Tecnai F20) was used to study film microstructure. Cross-sectional TEM samples were prepared by the lift-off method using a focusedion-beam (FIB) instrument. Film thicknesses were determined using energy-dispersive X-ray spectroscopy (EDX, Oxford INCA 350 apparatus connected to the Hitachi S-4800 SEM instrument). Thicknesses were calculated from Ir L-lines with the GMRFilm program24 assuming an iridium bulk density25 of 22.56 g/cm3. The uncertainty in thickness determination was estimated to be 5%. Film crystallinity and texture were studied by X-ray diffraction (XRD, PANalytical X’Pert Pro MPD) using the

Figure 1. Growth rate versus H2 pulse length at 250 °C. Films were grown with 1000 cycles. Other pulse and purge lengths were fixed at 1 s.

Next, we compared the growth rates of the O2 and O2 + H2 processes in the wide temperature range of 200−350 °C (Figure 2). At the lowest temperatures, both processes had low but almost equal growth rates of about 0.2 Å/cycle, which is likely due to the faster nucleation of the O2 + H2 process, as described below. At 250−350 °C, the O2 + H2 process had a 10−20% lower growth rate than the O2 process.

Figure 2. Growth rate versus deposition temperature. Films were grown with 1000 cycles using 1-s pulses and purges. 15236


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195 °C in the Ir(acac)3 + O3 + H2 process19 and at 200 °C in the Ru(EtCp)2 + O2 + H2 process.22 In ultrahigh-vacuum studies, hydrogen has been observed to desorb from single-crystal iridium surfaces mostly at 100 °C or below.28−32 This is also in accordance with the scheme recently proposed by Lu and Elam22 for the Ru(EtCp)2 + O2 + H2 process at 150−200 °C, where the surface is assumed to be bare ruthenium after the H2 pulse. Knapas and Ritala19 also concluded that, in the Ir(acac)3 + O3 + H2 process, the hydrogen does not remain on the surface, as no gaseous byproducts were formed during the following Ir(acac)3 pulse. This is in contrast to the Ir(acac)3 + O2 process, where combustion also occurs during the Ir(acac)3 pulse because of the adsorbed oxygen layer. Thus, it seems that H2 serves only to remove the adsorbed oxygen, in agreement with the observed decrease in the growth rate. Based on the discussion above, we propose a reaction scheme for the Ir(acac)3 + O2 + H2 process (Figure 4). It is apparent

The number of ALD cycles was varied from 100 to 2000 at a deposition temperature of 250 °C, revealing good linearity in both processes after an initial nucleation delay (Figure 3). The

Figure 3. Film thickness versus number of cycles at 250 °C using 2-s pulses and purges. Insets: SEM images of films grown with 100 cycles. Lines represent linear fits.

nucleation delays were estimated by fitting a line to the thickness values and calculating the intersection of the fitted line with the abscissa, yielding nucleation delays of 70 and 20 cycles for the O2 and O2 + H2 processes, respectively. SEM images of films grown with 100 cycles showed that the film grown with the O2 + H2 process was almost continuous, whereas the film grown with the O2 process consisted of smaller separate islands (Figure 3 insets). The thicknesses of the films, assuming the bulk density, were determined to be 2.2 and 1.2 nm for the O2 + H2 and O2 processes, respectively. The improved nucleation of the O2 + H2 process is an important finding, as shortening the nucleation delay has been one of the main goals and challenges in developing platinum-metal ALD processes. In previous reports on adding a H2 gas20,22 or plasma21 pulse to an O2-based ALD platinum-metal process, the hydrogen pulse has been found to lower the growth rate, in agreement with the present study. Mackus et al.20 tested the (MeCp)PtMe3 + O2 + H2 process and found the growth rate to decrease by 20% at 300 °C compared to the respective O2 process, whereas Kwon et al.21 saw a 75% decrease in the growth rate at 270 °C for the Ru(EtCp)2 + O2 gas + H2 plasma process. The much greater decrease of the ruthenium growth rate can be explained by the subsurface adsorption of oxygen, leading to a higher oxygen uptake that has only been found for ruthenium.7,27 Lu and Elam22 used the Ru(EtCp)2 + O2 + H2 process at low temperatures between 150 and 200 °C, whereas they found the respective O2 process to be ineffective at those temperatures. They demonstrated film growth only on platinum-metal substrate, however, severely limiting the applicability of the process, and the deposited films were not characterized. To understand the Ir(acac)3 + O2 + H2 process, two issues related to the H2 pulse need to be considered: What happens when the oxygen-covered surface is exposed to H2, and does the hydrogen stay on the surface to react with Ir(acac)3? H2 most likely removes the oxygen that is left on the surface after the O2 pulse. In the O2 process, this adsorbed oxygen combusts some of the ligands during the Ir(acac)3 pulse, which reduces steric hindrance between the adsorbed metal precursor molecules and further increases their adsorption density and finally film growth rate. In previous ALD studies, H2 has indeed been found to react with the surface oxygen, releasing water at

Figure 4. Schematic model of the O2 and O2 + H2 processes. Ir(acac)x refers to partially combusted precursor.

that there are three main differences in surface composition between the O2 and O2 + H2 processes. First, the surface that Ir(acac)3 adsorbs on is different, which changes the reactions that occur during and after the adsorption. The second difference follows from the first one, as oxygen causes the precursor ligands to combust partly in the O2 process, whereas on a bare iridium surface, the Ir(acac)3 molecules either remain intact or dissociate on the surface, but no byproducts are released. Third, in the O2 process, the surface is always covered with either oxygen or adsorbed precursor, whereas in the O2 + H2 process, the bare iridium surface is exposed between the H2 and Ir(acac)3 pulses. The second and third differences likely cause differences in surface diffusion, which can further modify film properties. Mackus et al.33 also speculated that differences in diffusion on metallic and oxidized surfaces play an important role in ALD of platinum. Film morphology was studied using SEM (Figure 5). The films grown using both processes had a clear granular morphology with a relatively small grain size. The O2 + H2 process appeared to produce films with slightly larger grains on the surface, which is somewhat surprising considering the lower 15237


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Figure 5. SEM images of iridium films. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges. Film thicknesses shown were determined by EDX.

growth rate and, hence, the lower film thickness compared to the O2 process. The grains appeared to grow larger with increasing deposition temperature, but the increase in grain size with increasing film thickness should also be taken into account (Figure 5b,e,f,h,k,l). AFM images revealed morphology trends similar to SEM while also yielding numerical information on film roughness (Figure 6). The O2 + H2 process mostly produced somewhat rougher films in comparison to the O2 process, although all of the films were relatively smooth for crystalline ALD films in general. At first, it seemed that temperature had almost no effect on film roughness (Figure 7a), but when the roughness was normalized with respect to film thickness, a trend showing a decrease in roughness with increasing deposition temperature was revealed (Figure 7b). This conclusion is supported by the fact that film roughness increased with increasing thickness at 250 °C. Roughness values determined by XRR for comparison were in good agreement with the AFM values, as the differences between the techniques were only 0.1−0.3 nm. AFM images were also used to estimate lateral grain size. As in the interpretation of the SEM images above, grain size was found to increase slightly with increasing deposition temper-

ature when considering 1000-cycle films (Figure 8). However, keeping in mind the increase in growth rate with increasing temperature, it seems that neither deposition temperature nor film thickness has had a strong effect on the grain size. Differences between the processes were small at 250 °C and lower temperatures, but at 300−350 °C, the O2 + H2 process produced films with larger grain size, which might be explained by the aforementioned differences in surface diffusion. Grain size estimation was also attempted using XRD, but this was found to be difficult because of the strong (111) texture of the films, which is discussed below. The TEM images in Figure 9 show that the films mostly consisted of tapered grains that were only slightly wider at the film surface than at the interface with the substrate. Although the majority of the grains extended throughout the film from the substrate to the surface, there were also some smaller grains with triangular cross sections that either ended or started in the middle of the film. The nearly equiaxed structures with lateral grain size of about 20−30 nm suggest that some of the original nuclei were eliminated in later stages of film growth, as large amounts of nuclei were observed after 100 cycles (cf. Figure 3 insets). According to the well-known zone models, the nearly 15238


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Figure 6. AFM images of iridium films. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges. Color scale shown in panel a corresponds to 25 nm in height.

equiaxed and columnar shape of the grains provides evidence for grain growth at the film−substrate interface even after film coalescence; this requires either bulk or grain-boundary diffusion of the film atoms, which is much slower than diffusion at the surface. As a result of surface diffusion only, the grains would adopt a strongly tapered shape instead, whereas without surface diffusion, very narrow fibrils would be formed. Some differences in surface morphology are also evident. Most of the grains in the film grown with the O2 process had sharp tops, whereas in the film grown with the O2 + H2 process, the sharp top is evident in some of the largest grains only. The greater roughness of the films grown with the O2 + H2 process (Figure 7) appears to have resulted from these largest grains, whereas the areas between them seem to be smooth. A fairly similar microstructure with almost equiaxed but less-sharptopped grains was also seen in TEM images of Ir films grown on Al2O3 using the Ir(acac)3 + O2 process at 300 °C.34 X-ray diffraction was used to characterize film crystallinity and texture. All of the films were found to be crystalline and oriented toward the [111] direction (Figure 10). Although the (220) reflection of iridium at 69° (2θ) was hidden by the (400) reflection of the single-crystalline silicon substrates (not

shown), the (200), (311), and (222) reflections of iridium (PDF 6-598) with varying but low intensities were also seen in all samples, in addition to the strong (111) reflection. To estimate the (111) texture, we calculated the intensity ratio of the two most intense unique reflections, (111) and (200), which, in a randomly oriented powder, is about 2 (PDF 6-598). It can be seen that the (111) texture generally increased with increasing deposition temperature (Figure 11). For the O2 + H2 process, this trend is clear across the whole temperature range, the intensity ratio reaching a value of 3000 at 350 °C, whereas for the O2 process, the (111) texture seemed to increase from 200 to 300 °C and then decrease from 300 to 350 °C. At all deposition temperatures, the films grown by the O2 + H2 process had larger (111)/(200) ratios than those grown by the O2 process, indicating a stronger texture. Rocking-curve measurements further confirmed the (111) texture with a full width at half-maximum (fwhm) of about 7° for both processes at the deposition temperature of 300 °C (Figure 12a). Trends in the rocking-curve fwhm (Figure 12b) were similar to those observed from the (111)/(200) intensity ratios, as a smaller rocking-curve fwhm corresponded to a larger intensity ratio. 15239


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Figure 9. Cross-sectional TEM images of films grown with (a) O2 and (b) O2 + H2 processes using 1000 cycles and 1-s pulses and purges at 300 °C.

Figure 7. (a) Absolute and (b) thickness-normalized film roughness versus deposition temperature. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges.

Figure 10. θ−2θ X-ray diffractograms measured from films grown at 250 °C using 1000 cycles and 1-s pulses and purges.

Figure 8. Lateral grain size versus deposition temperature as measured by AFM. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges.

To the best of our knowledge, such a strong (111) texture has not previously been reported for ALD iridium films. Even in rather thick films (up to about 70 nm), the (111)/(200) intensity ratio has been about 15 at highest,10,12,13 and no rocking-curve measurements have been shown. On the other hand, strong (111) textures have been observed for ALD platinum films: Aaltonen et al.35 measured a fwhm of 11° for the (111) rocking curve of a 110-nm film grown at 300 °C on Al2O3, and Potrepka et al.36 measured fwhm values of 6.84° for a 25-nm film and 4.58° for a 100-nm film grown at 270 °C on TiO2. Potrepka et al. also found the substrate to affect the texture of Pt films: A 25-nm film had (111) rocking-curve fwhm values of 6.84°, 6.92°, and 13.9° on TiO2, Al2O3, and HfO2, respectively. Some iridium films were also grown on soda lime glass substrates, and the (111) texture was observed to be much weaker than that of films grown on native oxide-terminated silicon. For example, a film grown with 1000 cycles at 300 °C using the O2 + H2 process had a (111)/(200) intensity ratio of

Figure 11. (111)/(200) peak intensity ratio versus growth temperature. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges.

900 on Si compared to only 14 on glass. Thus, we consider the substrate to be a possible explanation for the strong texture observed. Although both soda lime glass and silicon with a native oxide layer are amorphous on the surface, the effect of the crystalline silicon bulk versus the amorphous glass cannot be disregarded. Differences in chemical composition and homogeneity between the substrates can also cause changes in film nucleation, which could further affect the film texture. We propose a possible explanation for the observed texture differences between the two processes based on our reaction scheme (Figure 4): The bare iridium surface formed in the O2 + H2 process after the H2 pulse might allow faster surface diffusion, driving the crystals toward the favored (111) 15240


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Figure 12. (a) Ir(222) rocking curves measured from films grown at 300 °C and (b) fwhm of the rocking-curve peaks versus deposition temperature. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges.

Figure 13. Film resistivity versus (a) deposition temperature and (b) film thickness. In panel b, lines connect films grown at 250 °C to illustrate the effect of film thickness. Films were grown with 1000 cycles unless otherwise noted, using 1-s pulses and purges.

orientation more effectively than in the case of the O2 process, where the surface is always covered with either oxygen or partially combusted Ir(acac)3. It is also apparent that, because of the lower growth rate and longer cycle time, the iridium atoms have more time to diffuse in the O2 + H2 process, both on the surface and in the bulk of the film. However, we found that increasing the purge times to 5 s and thus allowing more time for diffusion did not strengthen the texture in either of the processes. Impurity contents were analyzed by TOF-ERDA, which can detect all of the elements starting from hydrogen. The results indicated that the films were very pure (Table 1). For both

temperature in both processes. It is well-known, however, that the resistivity of thin films increases strongly with decreasing thickness as a result of interface scattering,38 which can partially explain the observed trend. The effect of the deposition temperature, namely, the decrease of resistivity with increasing deposition temperature, was further confirmed by plotting resistivity as a function of film thickness and comparing films of nearly equal thicknesses (Figure 13b). There was also a difference between the two processes: Below 250 °C, the films grown with the O2 process had lower resistivities, whereas above 250 °C, the films grown with the O2 + H2 process had lower resistivities.

Table 1. Impurity Contents Measured by TOF-ERDA

CONCLUSIONS Iridium films were grown using two ALD processes: the commonly used Ir(acac)3 + O2 process and the Ir(acac)3 + O2 + H2 process developed in this study. The resulting films were then characterized. Our aim was to explore the effects of deposition chemistry on film growth and properties. Based on the deposition experiments and the literature we proposed a scheme explaining the O2 + H2 process. We believe that the H2 pulse removes the oxygen adsorbed on the iridium surface, forming water, which is followed by desorption of excess H2, finally resulting in a bare iridium surface. The addition of the H2 pulse expectedly decreased the growth rate by about 10− 20% depending on the deposition temperature, but interestingly, it also reduced nucleation delay, which will be a topic for our further studies. The film properties were studied for both processes and were found to depend on the applied process as well as the deposition temperature. Films grown by the O2 + H2 process were slightly rougher at all temperatures (200−350 °C) and had a slightly larger grain size at deposition temperatures above 250 °C, a lower resistivity at deposition temperatures above 250

deposition temperature 200 250 300 350 200 250 300 350

O (at. %) O2 + H2 Process 0.8 0.9 0.7 0.6 O2 Process 0.8 0.8 0.5 0.5

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C (at. %) 0.2 0.2 0.1 0.1 0.1 0.1 0.1 0.1

processes, the oxygen content decreased slightly with increasing temperature but was always less than 1 at. %, whereas the C content was always less than 0.2 at. % and the H content was less than the detection limit of ∼0.5 at. %. Both processes produced films with low resistivity (Figure 13a), approaching that of bulk iridium (4.7 μΩ cm at 0 °C).37 The resistivity seemed to decrease with increasing deposition 15241


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The Journal of Physical Chemistry C °C, and a stronger (111) texture at all deposition temperatures. All of the studied films were very pure, as their total impurity contents (H, C, O) were 1.5 at. % or less. We explained the differences in properties with changes in growth mechanisms depicted by the proposed reaction scheme, which, in turn, can lead to differences in diffusion processes. All of the films had strong (111) texture, which, to the best of our knowledge, had not previously been reported for ALD Ir films and would be an interesting topic for further study. We also believe that the obtained improvements over the O2 process, especially lowered resistivity and improved nucleation, warrant further interest in O2 + H2 processes, which seem viable for other platinum metals as well.

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(12) Kim, S.-W.; Kwon, S.-H.; Kwak, D.-K.; Kang, S.-W. Phase Control of Iridium and Iridium Oxide Thin Films in Atomic Layer Deposition. J. Appl. Phys. 2008, 103, 023517. (13) Hämäläinen, J.; Hatanpäa,̈ T.; Puukilainen, E.; Costelle, L.; Pilvi, T.; Ritala, M.; Leskelä, M. (MeCp)Ir(CHD) and Molecular Oxygen as Precursors in Atomic Layer Deposition of Iridium. J. Mater. Chem. 2010, 20, 7669−7675. (14) Hämäläinen, J.; Hatanpäa,̈ T.; Puukilainen, E.; Sajavaara, T.; Ritala, M.; Leskelä, M. Iridium Metal and Iridium Oxide Thin Films Grown by Atomic Layer Deposition at Low Temperatures. J. Mater. Chem. 2011, 21, 16488−16493. (15) Hämäläinen, J.; Puukilainen, E.; Kemell, M.; Costelle, L.; Ritala, M.; Leskelä, M. Atomic Layer Deposition of Iridium Thin Films by Consecutive Oxidation and Reduction Steps. Chem. Mater. 2009, 21, 4868−4872. (16) Kim, M. R.; Lee, J. H.; Choi, B. H. Effect of Hydrogen Plasma on Growth of Ir Thin Film by Plasma-Enhanced Hybrid Atomic Layer Deposition. Microelectron. Eng. 2012, 98, 400−404. (17) Kim, S.-W.; Kwon, S.-H.; Jeong, S.-J.; Park, J.-S.; Kang, S.-W. Improvement of Morphological Stability of PEALD-Iridium Thin Films by Adopting Two-Step Annealing Process. Electrochem. SolidState Lett. 2008, 11, H303−H305. (18) Kim, K.; Lee, S. Integration of Lead Zirconium Titanate Thin Films for High Density Ferroelectric Random Access Memory. J. Appl. Phys. 2006, 100, 051604. (19) Knapas, K.; Ritala, M. In Situ Reaction Mechanism Studies on Atomic Layer Deposition of Ir and IrO2 from Ir(acac)3. Chem. Mater. 2011, 23, 2766−2771. (20) Mackus, A. J. M.; Leick, N.; Baker, L.; Kessels, W. M. M. Catalytic Combustion and Dehydrogenation Reactions during Atomic Layer Deposition of Platinum. Chem. Mater. 2012, 24, 1752−1761. (21) Kwon, O.-K.; Kwon, S.-H.; Park, H.-S.; Kang, S.-W. PEALD of a Ruthenium Adhesion Layer for Copper Interconnects. J. Electrochem. Soc. 2004, 151, C753−C756. (22) Lu, J.; Elam, J. W. Low Temperature ABC-Type Ru Atomic Layer Deposition through Consecutive Dissociative Chemisorption, Combustion, and Reduction Steps. Chem. Mater. 2015, 27, 4950− 4956. (23) Nečas, D.; Klapetek, P. Gwyddion: An Open-Source Software for SPM Data Analysis. Cent. Eur. J. Phys. 2012, 10, 181−188. (24) Waldo, R. A. In Microbeam Analysis; Newbury, D. E., Ed.; San Francisco Press: San Francisco, 1988; pp 310−314. (25) CRC Handbook of Chemistry and Physics, 94th ed. [Online]; CRC Press: Boca Raton, FL, 2013−2014; http://hbcpnetbase.com (accessed Apr 20, 2014). (26) Jokinen, J.; Keinonen, J.; Tikkanen, P.; Kuronen, A.; Ahlgren, T.; Nordlund, K. Comparison of TOF-ERDA and Nuclear Resonance Reaction Techniques for Range Profile Measurements of keV Energy Implants. Nucl. Instrum. Methods Phys. Res., Sect. B 1996, 119, 533− 542. (27) Aaltonen, T.; Rahtu, A.; Ritala, M.; Leskelä, M. Reaction Mechanism Studies on Atomic Layer Deposition of Ruthenium and Platinum. Electrochem. Solid-State Lett. 2003, 6, C130−C133. (28) Chakarov, D. V.; Marinova, T. S. Interaction of Hydrogen with Ir(111) and Ir(110) Surfaces. Surf. Sci. 1988, 204, 147−160. (29) Lerch, D.; Klein, A.; Schmidt, A.; Müller, S.; Hammer, L.; Heinz, K.; Weinert, M. Unusual Adsorption Site of Hydrogen on the Unreconstructed Ir(100) Surface. Phys. Rev. B: Condens. Matter Mater. Phys. 2006, 73, 075430. (30) Hagedorn, C.; Weiss, M.; Weinberg, W. Dissociative Chemisorption of Hydrogen on Ir(111): Evidence for Terminal Site Adsorption. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 60, R14016−R14018. (31) Engstrom, J. R.; Tsai, W.; Weinberg, W. H. The Chemisorption of Hydrogen on the (111) and (110)-(1 × 2) Surfaces of Iridium and Platinum. J. Chem. Phys. 1987, 87, 3104−3119. (32) Liu, R.; Ehrlich, G. Chemisorption on Densely Packed Metal Surfaces: Mo, W, Re, and Ir. Surf. Sci. 1982, 119, 207−233.

AUTHOR INFORMATION

Corresponding Author

*Tel.:+358 2941 50193. E-mail: mikko.ritala@helsinki.fi. Author Contributions

All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Finnish Centre of Excellence in Atomic Layer Deposition.

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REFERENCES

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