Oxide Thin Film or Metallic Nanoparticles Formation - ACS Publications

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Buffer Layer Assisted Chemistry over Amorphous Solid Water: Oxide Thin Film or Metallic Nanoparticles Formation Liat Zilberberg, Shankar Harisingh, Serge Mitlin, Renana Elitzur, and Micha Asscher Langmuir, Just Accepted Manuscript • DOI: 10.1021/acs.langmuir.7b03203 • Publication Date (Web): 05 Feb 2018 Downloaded from http://pubs.acs.org on February 6, 2018

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Buffer Layer Assisted Chemistry over Amorphous Solid Water: Oxide Thin Film or Metallic Nanoparticles Formation L. Zilberberg, H. Shankar, S. Mitlin, R. Elitsur and M. Asscher* Institute of Chemistry, Edmund J. Safra Campus, The Hebrew University of Jerusalem, Jerusalem 91904, Israel

Abstract Novel procedures to grow pure thin metal oxide films are always welcome in view of their wide range of applications including photo-catalysis, solar cells, sensors and more. In this paper we present a unique way to grow pure nano films of metal oxides in-vacuum at the temperature range: 110--170 K. The Reactive Layer Assisted Deposition (RLAD) procedure for thin oxide films growth is based on the evaporation of a reactive metal element on top of a condensed layer of amorphous solid water (D2O-ASW). When applied to metals that do not react with the water layer the process yields metal nano clusters on the substrate. We observed that metal oxide films are formed if the redox potential is of 1.0 V or more, leading to deuterium molecules ejection to the gas phase (e.g. Ti and Al)while metals such as Zn, Fe and Ag, with redox potential less than 1.0 V, transform into nano-clusters, as revealed by SEM studies. We conclude that the redox potential ia a parameter that enables one to predict the nature and outcome of the ASW- buffer layer assisted chemistry.

Keywords: Metal-oxide, Reactive-Layer-Assisted Deposition (RLAD), Nanoparticles, Redox potential, Amorphous solid water (ASW).

*Corresponding author e-mail: [email protected]

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1. Introduction Thin films of metal oxides have widespread applications in many areas such as catalysis,1 coating,2 sensors,3 etc. There are various methods to prepare metal oxides such as Sol-gel,4,5,6 Atomic Layer Deposition (ALD),7 Chemical Vapor Deposition (CVD),8 Plasma spray9 and more. The Sol-gel process is probably the most frequently used, in particular for the coating of large areas by e.g. SiO2 and TiO2 near room temperature. The sol-gel process involves the hydrolysis of precursors such us metal-alkoxides or metal chlorides, namely the replacement of the OR group with a hydroxyl group. While the sol-gel procedure takes place at the liquid phase, an alternative technique is ALD which is a gas phase process enabling better conformity, thickness control and resolution. 7 Reactive Layer Assisted Deposition (RLAD) has been introduced as a way to grow very clean, thin metal oxide layers within ultra-high vacuum conditions. The RLAD method was first introduced for the preparation of MgO using condensed oxygen as the first reactant with Mg atoms evaporated on top

10

Water was later introduced as

the metal oxidizing agent for the preparation of TiO2 nano particles on Au (111) substrate11 and as a continuous thin film on various substrates12. Unlike the standard sol gel, the first step of this method is the oxidation of the metal element that takes place at 110K. Subsequently, condensation and aggregation occur simultaneously at low temperature (170K), as the buffer water molecules desorb. If the evaporated metal does not react with the condensed water molecules the final result is the formation of metal or oxide nano-particles deposited on the substrate during the buffer layer desorption. This procedure to prepare metal nano particle is known as buffer layer assisted growth (BLAG). In this procedure the size and the density of the nano clusters are controlled by the water buffer layer thickness and the amount of evaporated metal. Originally it was demonstrated with Xe as the buffer material for metallic nanoclusters growth.13,14 In a previous work we examined the formation of TiO2 employing RLAD, where a Sol-gel like mechanism was introduced, where Ti atoms reacted with the D2O-ASW film to create titanium oxide films. We found that the metal oxide morphology depends on the nature of the substrate, following high temperature annealing.12 2 ACS Paragon Plus Environment

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The RLAD approach enables one to prepare materials of high purity with a very good control on the film thickness. The mechanism is similar to the atomic layer deposition (ALD) method.7 However in the RLAD process the thickness can be controlled in a single step, while termination of the process occurs at the end of the evaporation cycles, all under UHV environment. Here we demonstrate the utilization of the RLAD procedure for the preparation of a new metal oxide, aluminum oxide, demonstrating extraordinary uniform and smooth film. We offer a generic scheme, enabling one to predict whether metal oxide films or nanoparticles will form with ASW as the buffer layer.

2. Experimental Samples were prepared under Ultra High Vacuum (UHV) environment at a base pressure of 1·10-10 Torr as described in previous publications.12,15 Different materials were used as substrates for various characterization purposes: 12×6×0.5 mm3 (n-type) SiO2/Si (100) wafer for SEM measurements and ultra-thin amorphous carbon (a-C) on a copper grid, the standard sample holder for transmission electron microscopy (TEM) inspections. In addition, 12×6×1.0 mm3 sapphire plates were used for optical absorption studies. Sample preparation procedure involved two steps after the substrate was cleaned (isopropanol sonication for 5 minutes, outside) and then inserted into the UHV chamber: The first step is the deposition of deuterated amorphous solid water (D2OASW) on a cooled (~110K) substrate followed by a direct deposition of the metal atoms. The substrate was clamped to a stain less steal plate, inserted to the vacuum via an introduction chamber and then cooled by attaching the sample holder at the bottom of a liquid nitrogen Dewar. D2O molecules were introduced by backfilling the UHV chamber and were adsorbed on the SiO2/Si (100), the substrate we used for SEM analysis. Layer thickness was determined by assuming unity sticking probability of the D2O molecules with exposure calibration such that 1 Langmuir (1L= 10-6 Torr·sec) equals 1ML. The cooled sample (110K) was exposed to 50-100L D2O in order to get 50-100ML D2O-ASW buffer material. At this temperature a compact ASW layer is formed.16 Using D2O improves the detection sensitivity of the 3 ACS Paragon Plus Environment

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increments of deuterium molecules released whenever pulses of metal atoms interact with the D2O-ASW film during the RLAD process. The second step involves the evaporation of the metal elements that are directed onto the water layer. Ti, Al and Zn were evaporated by using an e-beam evaporator (McAllister). Silver metal was evaporated by employing a resistively heated tungsten filament (0.25 mm in diameter) wrapped around a 1.0 mm in diameter, 5 mm long silver wire of 99.99% purity (Aldrich). In all cases, the evaporation rates were calibrated by employing an in-vacuum quartz micro balance (QMB) positioned where the sample is located (the dewar on which the sample is held can be Z-axis moved using a bellows system). The QMB units are typically deposition rates in Å/sec or total amount of deposited elements in Å. A calibration is needed for more realistic units in which 3Å is equivalent to 1monolayer (ML) that is about (5±2)∙1014 atoms/cm2, as a rough estimate. In the rest of this manuscript we'll use the QMB units (Å/sec) and in the proper places this value will be translated to the more physical value for evaporation rates of atoms/sec∙cm2.

In order to further verify the

evaporation rate, ex-situ profilometer measurements were performed. The profilometer determines the height of a material-step that was prepared by direct deposition of the metal following calibration by the QMB. For the Ti and Ag the evaporation rates extracted from the profilometer were in accord with QMB results, but for the softer metals namely aluminum and zinc we noticed a significant deviation. For these metals, the evaporation rate determined by the profilometer was twice faster than the rate measured by the QMB. For the sake of consistency the invacuum values obtained by the QMB were used throughout. The RLAD products, D2 molecules (when Ti and Al are evaporated on the ASW layer) were monitored by a quadrupole mass spectrometer (QMS) located at a fixed position about 5 cm from the sample. In order to improve the signal to noise ratio, the flux of the evaporated metals was modulated every 50-70 seconds by employing a magnetically coupled shutter between the metal source and the sample. After completion of a metal deposition step, the sample was heated to 300 K, such that the remaining adsorbed water molecules could react metal atoms that were not yet oxidized while desorbing. In order to initiate a crystallization process of the oxide film, annealing to higher temperatures (in-vacuum up to 1100 K, the maximum possible in our set-up) was carried out. Multilayer oxide films (titania and alumina) were obtained by performing multiple cycles of the water adsorption and metal 4 ACS Paragon Plus Environment

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evaporation on top. Our notation for the metal deposition cycles is as follows: 20 cycles of 5 Å Al on top of 50MLD2O on sapphire substrate will be: 5Å(Al)/50MLD2O/ x20/Sapphire.

The overall thickness of the amorphous oxide

films prepared by multi RLAD cycles were eventually determined by ex situ ellipsometry (J.A.Woollam alpha-SE) and profilometer (Veeco Dektak 150). The oxide films, as well as the metal nanoparticles grown on different substrates, were subsequently further examined by UV−vis, SEM, TEM. The chemical composition of different films and their oxidation state were determined by XPS (Kratos Axis). Optical properties (both absorption and light scattering utilizing the attached integrating sphere) were characterized by a UV−vis spectrophotometer (Cary 5000). Morphology was examined by the HR-SEM (Magellan TM 400L), AFM (Dimension 3100 Nanoscope V) and TEM (Tecnai F20 G2) instruments.

3. Results and Discussion In the course of these studies we have noticed that under similar experimental conditions some of the metals form oxide nanofilms whereas others are growing as metal nano-clusters. In the present work, we applied the RLAD method using different metals (Al and Zn), attempting to grow thin oxide films using ASW as the buffer material. By comparing these results with our previous reports on Ti (RLAD).12 Ag,15 Au (BLAG)17,18 thin film preparation, we introduced new insights into the thin film formation mechanism. The main objective of the present work has been to extend the buffer layer assisted deposition/growth methods on D2O-ASW while presenting a generic mechanism that will help one to predict whether an oxide film (RLAD) or metallic nano-clusters (BLAG) will form.

3.1 Ti and Al on D2O-ASW: Evolution of D2 molecules 3.1.1 In this section we will briefly describe the growth mechanism of TiO212,15 thin film via RLAD and compare it to the results of evaporation of Al on D2O-ASW. The 5 ACS Paragon Plus Environment

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impinging Ti atoms react with and become oxidized by D2O molecules while releasing D2 molecules to the gas phase, a unique signature of this mechanism. The ejected deuterium molecules were detected in real time during the modulated Ti atoms impingement by the QMS. When the magnetic shutter was open (typically for 50-70sec), the ASW was exposed to the Ti evaporator and the D2 pressure abruptly increased. A rapid decrease of the D2 signal occurred immediately after the shutter blocked the metal flux. This confirms the first step of the RLAD mechanism, namely, the monomer formation: Ti + D2O

Ti-OD + Dad and subsequently the two Deuterium radicals recombine

to form D2 that immediately desorb. In order to prepare a thicker layer, Ti was evaporated on ASW several times (e.g. 40 cycles) in a sandwich mode with only a single annealing step at the end, namely: 5ÅTi/50MLD2O/5ÅTi/50MLD2O/ X40. Such a process of multiple sandwich steps was previously demonstrated for the formation of large Ag nanoparticles.15 The thickness obtained following 40 sandwich cycles was measured by elipsometery to be 18±2 nm. The next oxidation step occurs while heating the sample to room temperature during desorption of the excess water molecules that take place at 150170K.11 The final step of oxidation is obtained by annealing to room temperature, resulting in (partial) crystallization of TiO2 particles.12

3.1.2 During the evaporation of aluminum on the D2O-ASW layer we noticed a similar behavior to that of the titania formation, namely, the D2 pressure increased as long as the water layer was exposed to the Al atoms flux. The D2 (𝑚𝑚⁄𝑧𝑧 = 4) partial

pressure measured by the mass-spectrometer during Al evaporation is shown in Figure 1 for three different D2O-ASW layer thicknesses.


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Figure 1:

Real time mass spectrometry detection of the evolution of D2 during Al deposition at a constant rate on D2O-ASW at the indicated ASW film thicknesses. The aluminum atoms impingement flux was 0.1 ±0.01 Å/sec and it was modulated every 65 sec. A flux of 0.1 Å/sec∙cm2 taken from the QMB is estimated to be (1.5±0.6)∙1013 atoms/sec∙cm2.

Whenever the magnetic shutter was at the open position, namely, the D2O-ASW was exposed to the evaporated Al atoms there was an abrupt increase in the partial pressure of D2. There is also a very small increase in the pressure of mass 3 (DH molecules, not shown). The formation of these molecules indicates the presence of H2O and some HDO molecules in the D2O reservoir. It seems that the amount of ejected D2 molecules does not change significantly for ASW layers equal or thicker than 50 ML. For 20 ML of D2O-ASW we can clearly see that the partial pressure of the emitted D2 is lower, possibly due to incomplete and not dense enough D2O-ASW layer. During the modulated cycles of Al deposition there is a gradual decrease in the rate (peak intensity) of D2 formation after the D2 ejection reaches a maximum. This behavior indicates that, as time elapses, the Al atoms gradually interact with less available D2O/OD groups that they can react with. Similar decay in the deuterium signal was previously obtained in the Ti /D2O-ASW system.12


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Al atoms do not penetrate deep into the ASW film but rather react with the topmost D2O layers. Therefore the gradual decrease of the ejected flux of D2 (at 250 seconds and beyond) depends on the flux and lateral density of the Al.

Figure 2: D2 pressure vs. time during the modulated evaporation of Ti (black), Al (red) and Zn (blue) metals. The impinging flux is identical for both metals (0.015 Å/sec∙cm2 or (2.5±1)∙1012 atoms/sec∙cm2. The blue line (for Zn) shows some modulated deuterium signal if amplified by a factor of 100.

In order to get a better resolution of the D2 temporal profile at the initial stages of the Al evaporation, we reduced the evaporation rate from 0.1 Å/sec (Figure 1) to 0.015 Å/sec. This slows down the reaction and the rate of the D2 release (Figure 2 red line) and makes it easier to follow the gradual increase of the rate of D2 formation. A comparison between the Ti and the Al reactions with the D2O ASW layers leads to the following conclusions with respect to the mechanism of alumina formation. Figure 2 presents the QMS measurements at mass 4 (D2) during the Ti evaporation (black line) and Al (red line), at identical metal deposition rates. The comparison reveals that the amplitude of the D2 pressure during the Ti evaporation is (initially) significantly


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higher than that observed during the Al evaporation. This is primarily attributed to a different deuterium atoms generation mechanism (see discussion below).

The oxidation state of Ti atoms (+4) in the TO2 is higher than that of Al (+3) in Al2O3, at a ratio of 1.3, if the metals are fully oxidized. Therefore, during the monomers formation step every Ti atom can react with up to four D2O molecules to form Ti(OD)4 compared to the reaction of Al with nominally three D2O molecules, resulting in Al(OD)3 formation. Consequently, one Ti atom can release a maximum of 4 deuterium atoms that immediately recombine and desorb at 110 K to form two D2 molecules. In the case of aluminum, each atom is responsible for a maximum ejection of 1.5 D2 molecules. Earlier work12 has suggested that the Ti atoms are not fully oxidized during the evaporation and their oxidation state at this step is on average +2. Therefore, relying on the experimental oxidation ratio, it can be concluded that one Al atom reacts with only one D2O molecule to form a single Al-OD group at this stage. Another fact that can explain the initial difference between the Al and Ti behavior is that the evaporation temperature of the Ti is much higher than that of Al. The evaporation temperature of a metal can be correlated with its melting point. The melting point of Ti is 1668 oC at ambient conditions and that of Al is only 660 oC.19 This enables the Ti atoms to penetrate somewhat deeper into the D2O-ASW layer compared to the Al atoms that cannot penetrate beyond the top few D2O-ASW monolayers. The large excess initial D2 formation in the case of Ti compared to Al is however primarily dictated by the different reaction mechanism of the metals with D2O, see below. We can explain the different time dependent D2 evolution peaks at the onset of the metal evaporation by exploring several possible mechanisms of the reaction.

1) 2) 3) 4)

𝑘𝑘1

𝑀𝑀 + 𝐷𝐷2 𝑂𝑂 → 𝑀𝑀𝑀𝑀𝑀𝑀 + 𝐷𝐷 ; M = Ti, Al 𝑘𝑘2

𝑀𝑀 + 𝐷𝐷2 𝑂𝑂 → 𝐷𝐷 − 𝑀𝑀 − 𝑂𝑂𝑂𝑂 𝑘𝑘3

𝑀𝑀 + 𝑀𝑀𝑀𝑀𝑀𝑀 → 𝑀𝑀𝑀𝑀𝑀𝑀 + 𝐷𝐷 𝑘𝑘4

𝐷𝐷 + 𝐷𝐷 → 𝐷𝐷2(𝑔𝑔)


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The first step is the reaction of a metal atom with water molecules, with a rate constant of k1, to produce the monomer M-OD and a physisorbed deuterium atom.20,21 The deuterium atoms formation rate is proportional to the product of the concentration of the reactants, Al atoms and water molecules, and the rate constant k1. A competing reaction should also be considered in which the metal atom condenses into the D-O-D bonds to form: D-M-OD with a rate constant k2. This reaction was suggested in the case of iron atoms interacting with D2O over FeO(111) 22. This reaction pathway does not contribute to the D atoms formation, therefore no D2 ejection to the gas phase upon metal atoms deposition is anticipated. The third step is the reaction of a metal atom with the already formed monomer M-OD to produce M-O-M and another physisorbed deuterium atom at a rate constant of k3. In the Al evaporation case, k3 is apparently larger than k1 and k2, the first step becomes rate determining for the deuterium atoms formation and initially only small amount of D2 ejection is observed. As more Al-OD species accumulate, the deuterium formation rate increases, eventually becoming similar to the deuterium formed by Ti atoms. The final reaction is the recombination of two physisorbed deuterium atoms to form one deuterium molecule (equation 4) that is immediately released to the gas phase. This step is very fast23 k4>> k1, k3, therefore the previous steps are the rate determining. In the case of the Ti atoms, k1 is significantly larger than that describing the reaction of Al with ASW. As a result, the D2 pressure increases instantly upon collision and reaction of the Ti atoms with the D2O-ASW film, as discussed above. At longer evaporation time the amount of the D2 molecules ejected to the gas phase gradually decreases. As previously mentioned, the reason for that is a gradual decrease in the density of available Al-OD species due to an increase of the metal coverage. It can be seen that at this stage (longer than 250 seconds) the behavior of the two metals becomes similar because there are enough Al-OD groups. Since the metals evaporation rate is the same for both metals, the final metal-oxide coverage is expected to be similar. The D2 pressure in Figure 2 does not decrease to zero within the measurement time (750 seconds) because we are dealing with porous amorphous solid water for which a full coverage to block all the OD groups requires a significantly longer evaporation time.


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The sample continues to oxidize during the annealing step to room temperature while the water molecules desorb (above 170K) and it is obviously further oxidized when the sample is exposed to ambient environment. Based on the AFM and TEM results (Figure 3), the morphology of 18±2 nm thick aluminum-oxide is a continuous and rather smooth (±0.25 nm roughness) thin film. Moreover, after in-vacuums annealing to 1100 K (the maximum temperature possible in our set-up) for 15 minutes on an amorphous carbon (a-C) substrate of the TEM sampled holder, the film is kept uniform and continuous but amorphous, as seen in Figure 3B and absence of any electron diffraction (not shown). This is probably due to the fact that not all the aluminum atoms are fully oxidized therefore crystalline phase cannot be obtained. As mentioned, the thickness of the oxide layer can be controlled by the number of repeated RLAD cycles. A sample prepared by 20 cycles of evaporation of 10 Å Al, determined in-situ by the quartz micro balance (QMB), on top of 50 ML ASW (10Å Al/50MLD2O/10ÅAl/50MLD2O/x20 ) was ex-situ examined with elipsometer and profilometer and found to have 18±2 nm thickness. Namely, evaporation of 1Å aluminum contributes about 1Å to the alumina thickness. This is very similar to the RLAD-TiO2 results.12 It is important to note that the

volume of the unit cell of Al2O3 (corundum) is almost twice larger than that of the TiO2 (rutile) in their crystalline phase. However, in our case we are dealing with the amorphous oxide with some uncertainty about the oxidation state, therefore the volume ratio is only an estimate.


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Figure 3: A) AFM image of a sample that was formed via the following procedure: 5Å (Al)/ 50ML (D O)x40 on SiO /Si (100) substrate. The surface roughness (±0.25nm ) was 2

2

measured along the (green) line (see insert). B) TEM image of a sample prepared via RLAD deposition procedure of 5Å(Al)/ 50ML (D O)/x3 (on amorphous carbon (aC) film on a2

copper grid) substrate after annealing to 1100K for 15min in UHV. EELS measurements show the presence of the Al and O elements. Considering instrumental sensitivity correction a 1:2 (Al:O) ratio is obtained (insert).

XPS measurement (Figure 4) of a sample that was prepared via RLAD of 5Å(Al)/ 50ML (D2O)/x20 supports the QMS results that the Al is fully oxidized by the interaction with water. This measurement reveals the two main peaks of alumina, one at 74.4 eV belongs to Al 2p

3/2

in Al2O3 and the second at 74.9 eV is assigned to

Al 2p 1/2 also in Al2O3. There is no evidence for the peak of metallic aluminum at 72.7 eV Al 2p

24 3/2.

One may claim that oxidation of the aluminum may have taken place

and completed during exposure of the sample to ambient environment. While it cannot be totally ruled out, the main evidence for full oxidation within the UHV chamber during the RLAD process is that following in-vacuum annealing to 1100K no signs of dewetting or particles formation could be observed. Together with the deuterium evolution we conclude that full oxidation was achieved during the RLAD step.


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Figure 4: XPS analysis of a sample prepared via RLAD-alumina that was formed by (5 Å (Al)/50 ML (D2O)x20, subsequently annealed to 1100K. The Al (2p3/2) peak at 74.4 eV and Al (2p1/2) peak at 74.9 eV belong to aluminum oxide compound.

3.2 RLAD-alumina as a protecting film for AgNPs. A question arises how the level of uniformity of the RLAD-alumina could be tested. We have previously shown that small (~4nm diameter) silver nanoparticles' (AgNPs) surface plasmon resonance (SPR) absorption at 460nm (when deposited on RLADTiO2 thin film), quenches within 50 minutes when exposed to ambient environment (see Fig. 5B, black squares)15. We also demonstrated that RLAD-TiO2 thin cover of these AgNPs significantly slows down the quenching (Fig. 5B, red circles). A similar effect of the RLAD-alumina coating of these small AgNPs is demonstrated as a way to examine the RLAD-alumina layer assuming that very uniform film will efficiently block the oxidation of the AgNPs which leads to their SPR quenching. The results are shown in (Fig. 5B, blue triangles), demonstrating a very effective blocking of the oxidation and quenching event, apparently blocking almost entirely any access of ambient oxygen that is believed to lead for the quenching of the SPR absorption In more details, small silver nano particles (~4nm) were grown (via BLAG deposition) on top of RLAD-alumina - 5Å(Ag)/50ML(D2O)x3/ and then covered by another layer of the RLAD-alumina (5Å(Al)/50ML(D2O)x3) Two effects of the sandwiched structure on the surface plasmon resonance (SPR) of the AgNPs were observed. The first relates to the plasmon resonance position (Figure 5A). In our previous work15 it was demonstrated that coating Ag NPs by a thin film of the RLAD titania have led to a major red shift (Fig. 5B, red line) compared to the absorption of bare silver nano particles on thin layer of TiO2 –(Fig. 5B, black line). The effect was attributed to the interaction between the metal oxide and the metal (Ag) nano particles that affects the hybrid material dielectric constant. In the present work we found that the surface plasmon resonance peak position of the embedded silver particles in alumina (blue line) is slightly blue shifted down to 405nm, compared to that of the bare silver (black line). It is important to mention that the substrate for the bare nano particles (black line in Figure 5A) was a thin RLAD- TiO2 which is responsible for the SPR absorption shift 13 ACS Paragon Plus Environment

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from 400 to 460nm.25 The alumina film does not lead to such a shift, probably as a result of the fact that this oxide has a wider band gap, therefore its insulating nature does not enable electrons exchange with the AgNPs. The second, and more important effect relates to the SPR stability (Figure 5B, blue triangles). In the case of alumina, the surface plasmon resonance is very stable compared to that of bare silver NPs. A decrease of only 10% in the plasmon resonance intensity after 60 min in air (blue triangles) for the silver particles embedded in alumina. In comparison, bare AgNPs plasmon intensity quenches completely in ambient environment during this period of time (Fig. 5B, black squares). This alumina/AgNP hybrid system is more stable even when compared to the AgNPs protected by TiO2 RLAD (red circles). As discussed above this may be attributed to the fact that the RLAD-alumina film forms a highly homogeneous and continues layer with very few structural defects that blocks ambient oxygen molecules almost completely from reaching the AgNPs, thus preventing its oxidation and SPR absorption signal quenching.

Figure 5: A. UV-VIS spectrum of bare ~4nm diameter AgNPs deposited on TiO2 covered sapphire (black squares), titania-covered AgNPs on TiO2 covered sapphire (red circles) and alumina-covered AgNPs on alumina covered sapphire (blue triangels). All spectra were taken within 2 minutes of removal of the sample from the UHV environment. Insets: A scheme of the embedded Ag nano particles in the metal oxide thin layer and the color explanation of the three SPR spectra. B. Decays of the AgNPs surface plasmon resonance intensity vs. time for the three samples: Bare AgNP (black squares), hybrid titania/AgNPs/titania (red circles) and hybrid alumina/AgNPs/alumina (red circles).


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3.3 Ag and Zn interacting with D2O-ASW The Ag evaporation was studied extensively in our previous article using the BLAG procedure.15 It was found that D2 molecules were not produced during the evaporation of Ag atoms on D2O-ASW indicating that these atoms (as expected, due to their relatively inert chemical nature) do not react with the D2O molecules. In order to get a wider view and better understanding of the RLAD vs. BLAG methods, Zinc atoms were studied here, since zinc oxide is useful in numerous applications and is a rather stable oxide material, however this element is also less reactive than early transition metal atoms. According to the QMS measurements the evaporation of Zinc atoms on top of the D2O-ASW layer has led to barely measurable D2 ejection at the same conditions used for Al and Ti in figure 2. Namely, similar to Ag, zinc atoms do not react effectively with the D2O molecules under these conditions. At higher metal atoms flux and QMS sensitivity one could see some signs of D2 ejection (less than two orders of magnitude smaller signal, as demonstrated in Fig. 2, blue signal). There are some works reporting that zinc reacts with steam.26 This may imply that the zinc oxide could be formed through evaporation at a suitable, higher temperature environment. Examination of the Zn samples formed on SiO2/Si (100) substrates utilizing XPS (Figure 6) and HR-SEM (Figure 7) support our conclusion that as a result of the ZnD2O interaction, Zn and/or Zn-oxide nanoparticles are formed. According to the XPS measurements (Figure 6) the electron binding energy (B.E.) of Zn (2p3/2) in zinc oxide at 1022.6 eV and the B.E. of neutral zinc atoms (2p1/2) at 1021.6 eV are very close to each other, practically impossible to distinguish between the two peaks.

Nevertheless, in this figure there is a small shoulder at the lower energy edge of the peak which indicates the contribution to this peak by some metallic zinc. We attempted to increase the intensity of the XPS peak at 1021.6 eV (the neutral Zn atoms) by the evaporation of more metal atoms during the BLAG/RLAD formation of Zn (or ZnO) on the SiO2 /Si(100) substrate. Nevertheless, XPS measurements performed on both high dose and lower dose Zn samples result in identical XPS 15 ACS Paragon Plus Environment

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spectrum, as shown in Figure 6, following a total dosage of 6Å Zn on the D2O-ASW layer (equivalent to 2±0.5ML Zn). Modified Auger signals (see reference 29 for Auger signal analysis) obtained from the low dosage Zn particles were monitored, see Figure 6B. The relevant Zn oxide (Zn+2) peak of the L3M4,5M4,5 transition at 2009.68 eV and the metallic (Zn0) peak at 2013.58 eV are clearly distinguishable. From this analysis metallic Zn0 contribution of 19 ±5 % was obtained. We note that area under the peak calculations from such Auger spectra contain high level of uncertainty.

There is another possibility that Zn interacts with D2O via reaction pathway number 2 to form D-Zn-OD in a significantly faster rate than reactions number 1 and 3. The SEM images and the QMS data have led us to propose that the majority of the zinc atoms react to form D-Zn-OD, as reported in the literature30, but at the same time a fraction of the atoms tend to form metallic (neutral) zinc via the standard BLAG procedure. It seems that in both cases nanometer size clusters are formed on the silica substrate. More oxidation of the metallic fraction of the clusters may take place upon in-vacuum heating to room temperature with further oxidation on the way from our vacuum to the XPS machine vacuum chamber.

A.


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18000

B. 16000

14000

ά = 2013.58

L3M4,5M4,5

Zn0

ά = 2009.68 Zn2+

12000

CPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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10000

L2M4,5M4,5

Zn0

8000

Zn0

Zn2+

6000 Zn0 4000

2000 980

985

990

995

1000

1005

1010

1015

1020

1025

1030

K. E. (eV)

Figure 6: A. XPS analysis of a sample that was formed via 5Å Zn/ 50ML(D2O)/x6. The Zn (2p3/2) peak at 1021.6 eV reflects the metallic zinc while the Zn (2p3/2) peak at 1022.6 eV reflects the zinc oxide. B. Modified Auger signals (see reference 29 for Auger signal analysis) obtained from the low dosage Zn particles. The relevant metallic-oxide Zn+2 peak of the L3M4,5M4,5 transition at 2009.68 eV and the metallic peak of Zn0 at 2013.58 eV are clearly distinguishable. From this analysis the metallic Zn contribution is 19±5%.

Figure 7: SEM images of Zn grown on D2O on a SiO2/Si (100) substrate: The amount of Zn evaporated is (A) 1 Å on 50ML D2O, (B) 5 Å on 50ML (D2O) and (C) 5 Å on 100ML (x12) times.


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SEM images of Zn grown on 50ML D2O on a SiO2/Si (100) substrate: The amount of Zn evaporated was 1Å(A) and (B) 5Å on 50ML D2O and (C) 5Å on 100ML D2O/x20 (in order to form larger clusters). 5Å is equivalent to 1.7ML Zn, which is roughly equivalent to (5±3)·1014 atoms/cm2.

SEM measurements (Figure 7) exhibit the familiar behavior of the BLAG method, namely, the nano particles morphology that can be manipulated by the ASW thickness and the amount of deposited metal.13,15,17 By increasing the amount of metal evaporated larger NPs are grown and the density increases. Evaporation of 1Å Zn ((1±0.4)·1014 atoms/cm2) on 50ML D2O leads to the formation of nano particles of 8±3 nm (diameter), while evaporation of 5Å Zn ((5±2)·1014 atoms/cm2) on the same layer of water results in average diameter for the particles of 11±6 nm. The nano particles density was (1.26±0.03)·1011 particles/cm2 in the first case and (2.73±0.03)·1011 particles/cm2 in the second case. We observed that the XPS measurements could not differentiate between the small and larger clusters, in both cases the XPS spectra looked as in Figure 6A. In addition to Ag and Zn, gold atoms behave similarly as was reported earlier by our group.17

3.4 RLAD or BLAG Examination of the interaction of various metallic elements with condensed ASW via the RLAD or BLAG schemes has led us to search for the dominant chemical property that governs the low temperature interaction of metal/D2O towards RLAD or BLAG (oxide film or metallic nanoparticles formation). The first parameter we examined if it may correlate with the RLAD or BLAG pathways was the melting point of the metal, dictating the kinetic energy of the impinging metal atoms. The melting points (at equilibrium ambient conditions) are Ti- 1668 oC, Al- 660 oC, Ag-962 oC, Au- 1064 oC and Zn- 419 oC.19 At a glance one observes that the trend of the melting point doesn’t correlate with the activity of the metals. Silver and gold have higher melting point than aluminum, yet, they don’t react with the D2O molecules while aluminum, with a significantly lower melting temperature, shows reactivity (D2 molecules ejection ) during its evaporation. The 18 ACS Paragon Plus Environment

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1

kinetic energy is 𝐸𝐸𝑘𝑘 = 2 𝑚𝑚𝑣𝑣 2 where m is the mass and v is the mean velocity of the

atoms. The mean velocity is given by the Maxwell Boltzmann distribution:

𝑣𝑣 = �

8𝑘𝑘𝑘𝑘 𝜋𝜋𝜋𝜋

where k is the Boltzmann's constant and T is the metal atoms temperature

at melting point. The kinetic energy of the Ti atoms during evaporation is about 20 kJ/mol and the kinetic energy of the Al atoms during evaporation is only 10 kJ/mol. 𝜃𝜃 =428 These numbers should be compared with the bond energy (D-O) (𝐷𝐷298𝐾𝐾(𝐷𝐷−𝑂𝑂)

kJ/mol)19 in the D2O molecule, the bond that the Ti atom has to break in order to form

the monomer Ti-OD + D radical. We conclude that the kinetic energy is not sufficient for overcoming the potential barrier of the reaction. For the aluminum reaction with D2O the situation is the same. Namely, the driving force for these reactions cannot arise from the kinetic energy of the evaporated metals.

The second parameter we examined was the redox potential of the metals. By comparing the oxidation potential of the metals we reveal a trend that can explain the observed activity, as shown in Figure 8. For convenience, we reversed the sign of the standard oxidation potential values in Figure 8. Al and Ti have the highest oxidation potential among the different metals. The Ti oxidation from neutral Ti0 atom to Ti+3

ETi = -1.63V and Ti0 atom to Ti+2 ETi

= -1.37V19, while the Al oxidation to Al+3

EAl = -1.66 V. Moreover, they have

the lowest ionization energy, 658.8 kJ/mol19 for the first ionization of Ti and 577.5 kJ/mol for that of Al. The oxidation potential of zinc from Zn(s) to Zn+2 is -0.76 V, for silver from Ag(s) to Ag+ is ( +0.8) V and for gold is (+1.52) V.19 The ionization energies of Zn, Ag and Au are: 906.4 kJ/mol, 731.0 kJ/mol and 890.1 kJ/mol respectively.19 The somewhat lower activity of Al in spite of its higher oxidation potential could be explained on the basis of kinetic energy arguments in the previous paragraph, Ti atoms may penetrate deeper into the ASW layer compared to Al. Another possibility for the enhanced D2 formation by the Ti atoms is that the rate constants of the reaction are different (eqs. 1 and 3). In the case of Ti k1 is larger than k3 as opposed to the Al case where it was demonstrated that k1 is smaller than k3. Larger k1 leads to faster D2 generation at the initial stages of exposure to the metal atoms, as in the case of the TiO2 formation. The larger k1 of Ti can also be


Langmuir

rationalized by the higher maximum oxidation number (+4) of Ti compared with Al (+3), as discussed in section 3.1.

Attempting to apply the above conclusions to predict what will be the nature of interaction of iron atoms with the water layer we examined the iron parameters. Its oxidation potential is in the region of the Zn atoms, EFe = -0.44 v, following the oxidation from Fe(s) to Fe + or Fe +3. Preliminary results reveal a negligible increase of masses 4(D2), 3(HD) and 2(H2) during the evaporation of Fe on the ASW. This supports our conclusion that the main parameter that directs the RLAD/BLAG methods is the redox potential that needs to be above 1.0eV in order to obtain an efficient RLAD process. As stated before, in the case of iron atoms, stable D-Fe-OD species were form upon interaction of Fe with D2O 22, a reaction that does not emit D2 molecules to the gas phase. On the other hand, one cannot rule out the possibility that oxide clusters (in addition to the metallic clusters) are formed as the D2O-ASW layer desorb.

Ti D2 pressure (Torr)(10-8)

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3

Al 2

1

Zn

0 -2.0

-1.5

-1.0

Fe -0.5

Ag 0.0

0.5

1.0

Au 1.5

Redox potential (V) Figure 8: Maximum measured pressure of the ejected D2 molecules during the evaporation of different metals vs. their redox potential.


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4. Conclusion The possibility to employ the RLAD method with D2O-ASW as the reactive layer for the preparation of various metal oxides films under clean UHV conditions has been examined and demonstrated. It was found that the metal elements that react to form the relevant metal-oxide are titanium and aluminum that possess the highest redox potential among the examined elements. The rate constant (k1) for the first step of the reaction 𝑀𝑀 + 𝐷𝐷2 𝑂𝑂 → 𝑀𝑀 − 𝑂𝑂𝑂𝑂 + 𝐷𝐷 that leads to fast ejection of D2, is faster for Ti

compared to that of Al, while Al atoms react faster with the intermediate Al-OD species (reaction number 3). Consequently, during the evaporation of Al, the D2 ejection rate reaches a maximum value at a later time compared to the Ti atoms. . In contrast silver and zinc atoms do not react with the ASW as indicated by the QMS measurements during the evaporation of these metals on the D2O-ASW showing no D2 formation. Hence, for these metals the D2O-ASW serves as a buffer layer that controls the preparation of metal nano particles, as in standard BLAG process13,17,18 Annealing to room temperature the Ag or Zn metals on the D2O-ASW film yields uniform distribution of bare AgNPs or ZnNPs, adherent to the substrate surface. In the case of Zn we cannot rule out the competing process of oxide clusters formation via reaction number 2 although at a significantly slower rate.

We speculate that other metal oxides can be produced under similar conditions to form thin, high purity oxide films. Both magnesium (redox potential of -2.37 V) and cerium (redox potential of -2.33 V),27 have higher oxidation potential than that of the Ti and Al, therefore are considered to be good candidates for oxide film formation via the RLAD scheme. It will be interesting to examine the reaction of Si atoms with the D2O-ASW to form thin film of silicon oxide, since the oxidation potential of Si is only -0.91 V19 which is higher than Zn but lower than that of Ti. Examination of this metal may set the limit value of the oxidation potential parameter that would result in oxide formation via RLAD.

The RLAD procedure is very useful for the investigation of metal oxides. It enables us to examine the effect of different parameters on the resulting oxide characteristics, such as defects and doping, and even to produce hybrid materials, in order to get material of improved environmental stability. For example, with the RLAD procedure


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we can prepare a combination of metal nano particles with a layer of metal oxide on top in order to protect nanoparticles against oxidation. This was demonstrated here via the stabilization of AgNPs plasmon absorption by RLAD-alumina and RLAD-titania.

Acknowledgement:

The support by the stuff of the Harvey M. Krueger Family Center for Nanoscience and Nanotechnology of the Hebrew University of Jerusalem is acknowledged. We would like specifically to thank to Drs. Inna Popov, Vitaly Gutkin and Itzik Shweky for their help with the sample characterization. This work was partially supported by the Israel Science Foundation (ISF), by the German Israel Foundation (GIF) and by a grant from Israel National Nanotechnology Initiative (INNI) via its FTA program.

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