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Photoactivated Molecular Layer Deposition through Iodo-Ene Coupling Chemistry Mie Lillethorup, David S. Bergsman, Tania E. Sandoval, and Stacey F. Bent Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b01780 • Publication Date (Web): 12 Nov 2017 Downloaded from http://pubs.acs.org on November 13, 2017

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Chemistry of Materials

Photoactivated Molecular Layer Deposition through Iodo-Ene Coupling Chemistry Mie Lillethorup, David S. Bergsman, Tania E. Sandoval, and Stacey F. Bent* Department of Chemical Engineering, Stanford University, Stanford, California 94305, United States.

Abstract This work introduces photoactivated molecular layer deposition (pMLD) as a route to deposit organic nanoscale polymer films with molecular-level control. Surface-tethered acrylate polymers are obtained through a radical step-growth polymerization where a diene and a diiodo monomer, ethylene glycol dimethacrylate (EGM) and 1,3-diiodopropane (DIP) respectively, are sequentially dosed in the vaporphase under pulsed UV-irradiation. pMLD occurs with a constant growth rate of 3.7 Å/cycle, and both monomers display self-limiting saturation. Films deposited by pMLD exhibit excellent stability in organic solvents. Furthermore, annealing studies with in situ X-ray photoelectron spectroscopy (XPS) reveals thermal stability up to 350 ˚C in vacuum. The mechanism behind pMLD of EGM and DIP is proposed based on detailed characterization of the polymer films by XPS and Fourier transform infrared spectroscopy, growth modeling, and comparison with control studies of pMLD involving monofunctional precursors. The coupling chemistry of pMLD presented herein provides future possibilities to create apolar linkages in the formation of nanoscale organic films. Introduction The ability to deposit nanoscale organic films with molecular-level control is of increasing importance for a range of applications for which precision in composition and architecture is crucial, e.g. for nanoelectronics1 and catalysis.2,3 Various highly controlled polymer deposition techniques have been developed within the last decades to follow demand for device and feature miniaturization. Among the most popular techniques are surface-initiated (SI) controlled radical polymerizations (CRPs)4 and

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layer-by-layer (LbL) assembly.5,6 These are both solution based techniques which can tailor physicochemical properties of surfaces and interfaces and enable the creation of complex polymeric surface films. Chemical vapor deposition (CVD) polymerization is, on the other hand, a solvent-free technique used to synthesize high purity, conformal polymer thin films from vapor phase monomers.7 Two commonly employed CVD polymerizations, initiated and oxidative chemical vapor deposition (iCVD and oCVD),8,9 have monomers and either an initiator or an oxidant simultaneously delivered to a cooled substrate through the vapor phase. A radical chain growth polymerization is observed for iCVD, whereas oCVD proceeds via radical combination of oxidized monomers (step-growth polymerization). SI-CRP, LbL assembly, and CVD polymerization are all used in the fabrication of nanomaterials, but other emerging techniques, such as molecular layer deposition (MLD), may provide advantages for depositing polymers with molecular control.10 Such high control is in demand as dimensions of devices approaches the atomic scale. Similar to CVD, MLD is a vapor-phase technique. Molecular control in MLD is achieved by sequentially dosing two or more bifunctional monomers, similar to its inorganic analogue atomic layer deposition (ALD). The monomers react with the substrate in a self-limiting fashion, which results in deposition of one molecular layer at a time and allows for deposition of highly conformal films with precise control over the film thickness. In its pure organic form, MLD has been used to deposit polyimide,11,12 polyurethane,13 polyurea,14,15 polyamide,16,17 polyazomethine,18 polythiourea,19 and polyester materials.20 Within many of these polymer classes, a range of aliphatic and aromatic backbones has been employed to tune the properties of the final polymer film.14,21,22 Notably, the coupling chemistry utilized in each of these seven MLD systems all rely on the reaction between a good electrophile (acid anhydride, acid chloride, aldehyde, isocyanate, or isothiocyanate) and a strong nucleophile (amine or alcohol) to form a covalent, polar linkage. The reason why MLD is generally 2 ACS Paragon Plus Environment

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limited to this class of reactions is due to the vapor phase nature of MLD, which complicates use of typical organic chemistry tools such as tuning of catalysts, solvents, and pH. In fact, the only tunable parameters in traditional MLD are the temperature, the monomer concentration (pressure), and the reaction time. Thus, only thermodynamically and kinetically favored reactions will proceed and result in controlled MLD growth. This has greatly limited the development of new MLD polymers. Recently, the Parsons group demonstrated oxidative MLD (oMLD) using molybdenum(V) chloride to facilitate the growth of conducting poly(3,4-ethylenedioxythiophene) (PEDOT).23,24 This process is very similar to the oCVD, but it doses the monomer and oxidant sequentially. PEDOT is so far the only polymer deposited by oMLD. In a solution based process, known as molecular LbL, UV-light has also been used to facilitate molecular-level control of polymer growth.25 With dithiol and diene monomers, the controlled deposition of multilayer films was achieved through the photoactivated thiol-ene reaction. Herein we introduce all vapor-phase, photoactivated MLD where a direct energy input from UVlight is used to induce formation of covalent, nonpolar C–C linkages. Such an approach will expand the scope of available reactions applicable to MLD and thereby the range of polymers. An interesting class of materials in this context is organoiodine compounds with photolabile C–I bonds prone to undergo homolytic scission upon UV-irradiation.26,27 This facile bond cleavage was utilized by Wolpers et al. who demonstrated highly controlled UV-initiated iodine transfer polymerization (ITP) of n-butyl methacrylate.28 ITP is a solution-based CRP, where I˙ is reversibly transferred between iodo-capped polymers (Polymer–I) and propagating polymer chains (Polymer˙) to retain control.29 In another solution-based process, the chain transfer of ITP has also been used to facilitate radical step-growth polymerization of diene and diiodo monomers,30,31 very similar to the thiol-ene polymerization.



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In this work, we demonstrate photoactivated MLD (pMLD) from vapor phase precursors utilizing the radical step-growth polymerization of the diene and diiodo monomers, ethylene glycol dimethacrylate (EGM) and 1,3-diiodopropane (DIP). The surface-tethered acrylate polymers are interesting e.g. for patterning applications using electron-beam lithography. The monomers were chosen to serve as a model system for the pMLD process, but the same coupling chemistry may be expected to enable growth of hydrophobic and purely aliphatic polymers. Saturation behavior of both monomers and a constant growth rate confirm the controlled nature of the process. The chemical composition of the surface polymers as determined by Fourier transform infrared (FTIR) spectroscopy and X-ray photoelectron spectroscopy (XPS) agree with the coupling product of EGM and DIP monomers. We additionally show that these polymer films possess excellent thermal stability. A mechanism for pMLD is proposed based on the growth behavior of control pMLD studies involving monofunctional precursors and detailed XPS analysis. This pMLD process, using light to facilitate coupling reactions, introduces a route to create carbon-carbon bonds without the need for a catalyst, which is not achievable by existing MLD chemistries, and the pMLD process is thus an important tool for formation of new families of polymer thin films created with molecular-level precision. Materials and Methods Ethylene glycol dimethacrylate (EGM), 1,3-diiodopropane (DIP), 1-iodopropane (IP), methyl methacrylate (MMA), 98% sulfuric acid, and 30% hydrogen peroxide were purchased from SigmaAldrich, whereas 96% ethanol and HPLC grade acetone were purchased from Fisher Chemical. DIP and IP were stored in a nitrogen-purged glovebox until used, and EGM and MMA were stored in a refrigerator. As received, EGM and MMA reagents contain monomethyl ether hydroquinone (a radical inhibitor). Before use, the inhibitor was removed by passing the monomer through an Al2O3 column. Films were deposited on (100) silicon wafers with a 1.6 nm native oxide (measured by ellipsometry) 4 ACS Paragon Plus Environment

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purchased from WRS Materials. Batches of wafers were initially cleaned by piranha treatment, by immersing in a 7:3 ratio solution of sulfuric acid and hydrogen peroxide for 15 min before rinsing and storing in deionized water. (Caution: fresh piranha solution is hot and extremely corrosive. Proper training and extreme care should be taken when preparing and handling piranha solution.) Right before loading a wafer sample into the MLD reactor, the sample was dried and then cleaned for 15 min in a Novascan PSD Series Digital UV Ozone System to remove any remaining organic contaminants. MLD films were deposited using a hot-wall flow reactor pumped by an Alcatel rotary vane pump with a base pressure near 1 mTorr. An exterior UV deuterium lamp (Hamamatsu, L7292) was attached on top of the main chamber and centered to irradiate samples inside the chamber after the light passed through a long-pass filter (225 nm from Newport, 250 nm from Asahi Spectra, or 280 nm from Thorlabs) and a quartz (DUV fused silica, Corning HPFS 7980, transmission ≥99.8% at 248 nm and >80% at 200 nm) viewport from Kurt J. Lesker. The light intensity is estimated to be on the order of 0.1 µW/cm2, based on the specifications from the manufacturer. Irradiation of the samples was controlled through a home-built, automated shutter, positioned between the long-pass filter and the UV lamp. The reaction chamber and the monomer precursors were kept at room temperature in all experiments. A detailed description of the deposition conditions are included in the Supporting Information, but in brief, one MLD cycle included: 1) a dose of EGM (or MMA), 2) a soak time where EGM (or MMA) stayed in the reactor with continuous irradiation, 3) a nitrogen purge step, 4) a dose of DIP (or IP), 5) another soak time with continuous irradiation, and 6) another nitrogen purge step. A nitrogen purge optimization study is provided in Figure S1. Immediately after deposition, film thicknesses were measured by variable angle spectroscopic ellipsometry using a J. A. Woollam Co. α-SE spectroscopic ellipsometer with a spectral range of



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300−900 nm. Measurements were taken at incidence angles of 65 and 70° relative to the surface normal. The film thickness and the refractive index (nx) were modeled by employing a Cauchy model with zero extinction term. nx decreased with increasing film thickness, and generally fell within the range 1.8−1.57 going from thinnest to thickest films, a phenomenon well-known for thin films.32 The model and the non-fixed nx were chosen because they accurately modeled film thickness, as verified by X-ray reflectivity (XRR) measurements. The reported thickness values are the average value from measuring the thickness in three spots on each sample. FTIR spectra were taken using a VERTEX 70 FTIR spectrometer from Bruker with a germanium attenuated total reflectance (ATR) plate. Spectra were recorded at 4 cm−1 resolution with 200 scans, and a UV-ozone treated Si substrate was used as a background. XPS was performed on a PHI VersaProbe III spectrometer with monochromatic Al Kα X-ray source. The X-ray beam diameter was 200 µm with 50 W power. For survey spectra, 5 scans were recorded using an energy step of 0.8 eV and 224 eV pass energy, while high resolution spectra were recorded with 10 scans using an energy step of 0.1 eV and 55 eV pass energy. The spectra, calculated atomic percentages, and component ratios presented in this paper all originate from samples heated in situ to 75 ˚C, because it was observed that adventitious carbon and other sample contamination were removed by this annealing. Atomic percentages were determined from survey spectra following Shirley background subtraction. Deconvolution of high resolution spectra were fitted while constraining the full width half maximum (FWHM) from the different species to be identical. All spectra were calibrated against the C–H/C–C peak in the C1s spectrum (284.9 eV). The in situ thermal stability of pMLD films was tested by incrementing the temperature of the XPS heat stage in steps of 10 ˚C while letting the pressure of the XPS chamber stabilize for 5 min at each temperature step. XPS was recorded


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at 75, 100, 150, 200, 250, 300, 350, and 400 ˚C, after waiting 10 min at that temperature to ensure uniform heating of the substrate. XPS recording took another 10 min, so the heating stage was at each of these target temperatures for 20 min in total. XRR and X-ray diffraction (XRD) were performed on a PAN Analytical X'pert Materials Research Diffractometer with Cu Kα radiation (λ=1.54 Å) and generator power of 45 kV and 40 mA. For XRR a divergence slit of 1/32° and a parallel plate collimator receiving slit were used. 2θ-ω scans were taken with a step size of 0.005° and a scan time of 0.5°/step, after sample alignment. The data was then fit using PAN Analytical X-ray reflectivity software, assuming a Si/Interface/MLD film model. For the XRD, a divergence slit of 1/2° was used at the incidence beam (no parallel plate collimator receiving slit). Results To achieve pMLD on a silicon substrate, diene and diiodo monomers, EGM and DIP (Scheme 1), were dosed in the vapor phase in an alternating manner separated by nitrogen purge steps. In between dosing and nitrogen purging, each the monomer vapor was allowed to stay in the reactor for a period of time with continuous irradiation (soak time). Soaking steps are commonly used in ALD33 and in particular MLD11,14,20 to achieve saturation behavior. The growth of the layer-by-layer surface polymerization was followed by measuring the film thickness ex situ by ellipsometry. Homogeneous films were obtained, with thickness variations across the film generally < 0.3 Å and roughness measured by AFM ~1% of the film thickness (Figure S2). We initially tested the wavelength-dependency of the iodo-ene pMLD process, by passing the UV-light from a deuterium lamp through a longpass filter with a cut-on wavelength of 225, 250, and 280 nm, respectively. As seen from Figure 1, only minor pMLD was achieved for 16 MLD cycles (16×[EGM/DIP]) at wavelengths above 280 nm, whereas successful growth was observed in the case of 225 and 250 nm longpass filters, with the film resulting from 7 ACS Paragon Plus Environment


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irradiation down to 225 nm being significantly thicker than the one prepared with UV light > 250 nm. The n→σ* transition (a lone pair electron excited to the antibonding σ* orbital) of the DIP C–I bond is observed at 256 nm (Figure S3). The fact that only minor pMLD is achieved when using the 280 nm longpass filter indicates that the photo-induced homolysis of the C–I bond is involved in the pMLD process. The small pMLD growth with a 280 nm long-pass filter can be explained by some optical transmission through the filter down to ~260 nm, in combination with a broad absorption band for the n→σ* transition of the C–I bond. In a dark control experiment, no EGM/DIP deposition was observed. Scheme 1. Schematic Illustration of Molecular Layer Deposition. One cycle includes (a) exposure to ethylene glycol dimethacrylate (EGM) under irradiation followed by a nitrogen purge step without irradiation and (b) exposure to diiodopropane (DIP) under irradiation followed by a nitrogen purge step without irradiation. This cycle is repeated (c+d) to achieve polymeric surface films.

Interestingly, pMLD control experiments of 16 cycles EGM alone result in 21 Å surface film using the 225 nm longpass filter, whereas the film thickness was much smaller with the 250 nm filter (8 Å) and insignificant at 280 nm (3 Å, corresponding to minor physisorption of EGM or contaminants in the reactor). The UV-absorption spectrum of gaseous methyl methacrylate (MMA) shows an n→π* transition at 240 nm.34 Based on the structural similarities of MMA and EGM (the two methacrylate units in EGM are separated by a fully saturated ethylene group, and hence, the excitation of one ester group is not expected to be significantly influenced by the other intramolecular ester group), we expect that EGM has a very similar transition, and hence undergoes photopolymerization when using the 225


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nm longpass filter. A small overlap of the irradiation passing the 250 nm longpass filter with the n→π* transition of EGM may result in some photopolymerization of EGM and explain the observed 8 Å surface films. However, it is worth noting that EGM photopolymerization is self-limiting (up to 32 cycles of EGM does not add additional film thickness), and that the combined EGM/DIP pMLD process gives 6–7 times higher film thickness than EGM alone. Hence, the photopolymerization of EGM is only a minor contribution in the combined EGM/DIP pMLD process. Therefore, we continued our studies on pMLD of EGM/DIP using the 250 nm longpass filter.

Figure 1. Wavelength-dependent analysis of pMLD of EGM/DIP. The thickness of 16×[EGM/DIP] and 16×[EGM] are measured as a function of cut-on wavelength of the longpass filters.

In a true MLD process, each of the precursors reacts with the surface or surface adsorbed molecules in a self-limiting, monolayer-by-monolayer fashion.10 Herein, the saturation behavior of DIP and EGM was studied by varying the soak time of each of the two precursors separately in 16 cycles of EGM/DIP (Figure 2). As seen from Figure 2a, a rapid increase in thickness is observed when increasing the soak time of EGM. Saturation is achieved after 600 s, and no further growth is seen at higher EGM soak times. The saturation of DIP is slower, and is only seen after 3400 s (Figure 2b). However, at even longer DIP soak times, the thickness remains the same. It is furthermore noted that if no co-reactant is dosed, i.e. 16 cycles pMLD of EGM or DIP, only minor deposition is seen. The thickness is ~8 Å and



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10 Å for EGM alone and DIP alone, respectively. The self-limiting saturation of both precursors, and the fact that both EGM and DIP are required to achieve film deposition, suggest that pMLD of EGM/DIP is indeed a true controlled MLD process. The long processing time may be reduced by changing the UV-light source to a higher intensity lamp (e.g. mercury), optimizing chamber dimensions, or adjusting parameters such as temperature and flow rate of purge gas. Taking the long irradiation times into account, especially for DIP, it is apparent that no UV- or radical induced degradation occurs even at long soak/irradiation times, where the films thickness does reach a plateau (Figure 2b).

Figure 2. Saturation curves derived from 16×[EGM/DIP]. The thickness is measured as a function of (a) EGM soak time (DIP soak time = 3400 s) and (b) DIP soak time (EGM soak time = 600 s). In both cases the sample is UV-irradiated both during EGM and DIP soaking. The error bars on the 600 s/3400 s deposition represent


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the standard deviation (± 6 Å) from 4 independent runs, and reflect some variation between individual depositions. The MLD character of this photoactivated iodo-ene process was further confirmed by measuring the film thickness as a function of the number of MLD cycles (Figure 3) using the optimized soak condition (600 and 3400 s for EGM and DIP). A linear increase is observed up to 30 cycles, with a constant growth rate of 3.7 ± 0.3 Å/cycle (R2 = 0.95). This value is considerably lower than the extended length of one EGM-DIP repeating unit, 14.2 Å (obtained from energy minimization of the structure by density functional theory calculation (DFT), see Supporting Information). Such a trend is commonly seen for MLD of organic polymers.14,17 It was recently shown that termination events and not growth angle are more likely the primary driving force causing this difference.21 Taking the flexible nature of the EGM-DIP backbone into account, coupling between neighboring polymer chains by double reaction of one of the precursors may be expected. Notably, the growth curve intercepts the y-axis very close to the origin. It indicates that iodo-ene pMLD grown from blank SiO2 substrates does not have a nucleation delay period with lower growth rate, as is often observed for ALD.35,36

Figure 3. Growth curve for pMLD of EGM/DIP (black), where the black, dashed curved represents the linear fit to the experimental data points. pMLD using a monofunctional iodo precursor (EGM/IP, red) and a monofunctional ene (MMA/DIP, blue) is also included in the figure. The red, dashed curve is the best fit of Eq. 1 to the experimental thicknesses of EGM/IP (vide infra).



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To gain more insight into the mechanism of photoactivated iodo-ene MLD, we investigated two pMLD processes using (1) a monofunctional iodo precursor, 1-iodopropane (IP), in combination with EGM, and (2) a monofunctional ene-precursor, methyl methacrylate (MMA), in combination with DIP. We interestingly found that for (1), the initial growth rate is close to pMLD of EGM/DIP (Figure 3). However, with increasing number of cycles, the growth rate decreases as seen from the non-linear growth curve depicted in Figure 3. The growth rate does not go to zero, but rather approaches a steady state growth rate at high number of cycles (0.8 Å/cycle). In the case of (2), very limited growth is observed even after 30 MMA/DIP cycles. In general, it may be expected that any ideal MLD process involving a monofunctional precursor will lead to immediate chain termination, and hence no effective MLD growth. This was indeed seen for pMLD of MMA/DIP, where the minor growth is ascribed to photopolymerization of MMA, as discussed earlier.37 The fact that deposition is seen for EGM/IP suggests that the mechanism of iodo-ene pMLD is not as simple as that proposed in Scheme 1. This will be addressed in the discussion section below. The film density of the surface-tethered EGM/DIP films was found from XRR analysis (Figure 4a) to be 1.23 g cm-3. This value is very close to the known density of bulk PMMA, 1.17−1.20 g cm-3,38 which from its structural resemblance to the EGM/DIP polymer may be expected to have a similar density. Film thicknesses estimated from XRR are also in agreement with the values derived from ellipsometry. For example, the XRR thickness of EGM/DIP analyzed in Figure 4a is 127 Å, whereas the ellipsometric thickness is 118 ± 7 Å. The difference in these thicknesses is presumably due to uncertainties in the fitting of ellipsometry data. Interestingly, the EGM/DIP film shows some degree of ordering, as seen from the XRD pattern displayed in the Supporting Information, Figure S4. The main peak corresponds to a d-spacing of 5.2 Å, while the two smaller diffraction peaks have d-spacing of


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roughly 4.8 and 6.3 Å. The full width half-maximum values range from 0.2 to 0.6 Å. Previously, outof-plane ordering has been observed for polyamide39 and polyurea21 films grown by MLD. More advanced X-ray diffraction analysis of the EGM/DIP films is necessary to elucidate the origin of the Xray scattering reflexes observed here, but we speculate that they arise from backbone alignment of neighboring polymer chains. The chemical structure of the EGM/DIP surface films and the linkages formed in the pMLD process were analyzed by FTIR spectroscopy (Figure 4b). In the region 1850−1050 cm-1, the band at 1735 cm-1 is assigned to the C=O stretch of the unconjugated ester, the features at 1480−1235 cm-1 to C−H deformations from CH2 and CH3 groups, and the peak at 1167 cm-1 to the C−O stretch. The full conversion of the EGM monomer is confirmed by the absence of bands pertaining to the C=C stretch around 1640−1590 cm-1, and by the position of the C=O stretch band (1735 cm-1), which is blue shifted by ~10−20 cm-1 with respect to the corresponding conjugated ester of the EGM monomer (See also DFT calculated spectra of EGM and EGM-DIP and the experimental spectrum of EGM monomer in Figure S5a).40,41 In the C−H-stretching region around 3000 cm-1 displayed in the inset of Figure 4b, the bands at 2988, 2953–2933, and 2858 cm-1 are assigned to νa(O–CH2); νa(CH3 and CH2) and νs(O– CH2); and νs(CH3 and CH2), where νa and νs denote antisymmetric and symmetric stretches, respectively. The assignment is based on a DFT calculated spectrum of an EGM-DIP unit as shown in Figure S5b. It is noted that only a shoulder appears above 3000 cm-1, where C−H-stretching from sp2 hybridized centers is expected. By comparing to the DFT calculated spectrum, this shoulder is assigned to the C−H-stretching of O–CH2. This again confirms the complete radical addition to the C=C bonds of EGM, successfully converting these to sp3 hybridized centers.



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Figure 4. (a) XRR analysis of an EGM/DIP sample with 118 Å ellipsometric thickness and (b) representative FTIR spectrum of an EGM/DIP film deposited by 30 cycles pMLD.

Furthermore, the chemical composition of the EGM/DIP films was analyzed by XPS, where C and O were the main components detected, with traces of I (0.3%) also present. The C/O ratio is 2.7:1, whereas the theoretical value obtained from the expected EGM-DIP structure is 3.25:1 (Table 1). This ratio suggests that the surface film has a higher fraction of O than expected. On the other hand, compared to a scenario in which EGM is the only component of the film (in the case of EGM photopolymerization which would have a C/O ratio of 2.5:1), the actual surface film would be low in O. It is interesting to note that the composition of the film is more similar to EGM than EGM-DIP, despite the fact that pMLD of EGM alone does not result in any significant deposition. This will be elaborated on in the discussion section below.


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Table 1. C/O Ratio and Ratio between Components in High Resolution C1s XP Spectrum. The Theoretical Ratios for [EGM-DIP] and [EGM] Polymers are Included.

Component ratio by XPSa

-[EGM-DIP]- Theoretical

-[EGM]- Theoretical

2.7:1

3.25:1

2.5:1

C−H and C−C

2.49 ± 0.05

3.5

2

C*−(C=O)−O

1.02 ± 0.02

1

1

C−O

1.02 ± 0.02

1

1

C=O

1.00

1

1

C1s C/O

a

The thickness of the EGM/DIP film is greater than the XPS analysis depth, so no signal from the SiO2 substrate is observed.

From high resolution C1s and O1s XPS, detailed information on the different bondings in the polymer film is derived. Deconvolution of the C1s spectrum shown in Figure 5a gives four peaks at 284.8, 285.7, 286.8, and 288.9 eV, which are assigned to C−H/C−C, C*−(C=O)−O, C−O, and C=O,42 respectively, in agreement with the expected functional groups of EGM-DIP. The O1s spectrum (Figure 5b) shows two distinct components at 532.1 and 533.6 eV. The former corresponds to C=O groups, and the higher binding energy component is assigned to the more electron deficient C−O groups. These two components are also found in the structurally similar PMMA polymer,42 and the 1:1 ratio confirms that the MLD films contains binding modes in agreement with the expected EGM/DIP polymer.



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Figure 5. High resolution (a) C1s and (b) O1s XPS of EGM/DIP. Deconvolution into the relevant components is shown. The open circles represent the raw data points.

The ratio between the four C1s components is given in Table 1 together with the theoretical values for polymers of EGM/DIP and EGM, respectively. Interestingly, the fraction of C−H/C−C bonds is lower than expected for the ideal EGM-DIP repeating unit, but higher than expected for pure EGM. The same was true for the C/O ratio, being in between that of EGM-DIP and EGM. Both of these observations indicate that the MLD film contains fewer alkyl fragments than an EGM-DIP polymer, and suggests that more than one EGM monomer is added per DIP monomer in the iodo-ene pMLD process. This will be discussed in more detail below, with a discussion of the photoactivated iodo-ene mechanism. The C=O/C−O/C*−(C=O)−O ratio is in good agreement with the expected 1:1:1 ratio (Table 1).


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The stability of EGM-DIP films deposited by iodo-ene pMLD proved to be good. Upon storage at ambient lab conditions for 2 months, the C1s XP and FT-IR spectra of pMLD films were unchanged, while the thickness increased by 5%. This suggests that the polymer chains are themselves chemically stable but the film itself may change slightly, perhaps through the absorption of water or through a slight relaxation in the chain configurations. The stability was tested by sonicating samples in water, ethanol, and acetone. The thickness of a range of samples (20–70 Å) remained unchanged after 10 min sonication in any of these solvents with no notable delamination occurring. After 1 h sonication in water, up to 40% decrease in film thickness was observed on some samples, likely due to slow hydrolysis of the ester bonds in the polymer backbone. After 1 h sonication in acetone, a solvent expected to swell the pMLD film based on its chemical similarity to PMMA, the film thickness decreased less than 10%, demonstrating the high stability of the coating in organic solvents. The thermal stability of an MLD film ~190 Å in thickness was tested under vacuum in an in situ XPS heating experiment. The C1s and O1s high resolution spectra going from 25 ˚C to 400 ˚C are shown in Figure 6a and b, where the relevant component positions assigned in Figure 5 are indicated with dashed lines. When increasing the temperature from 25 to 75 ˚C, the pressure in the XPS chamber increased significantly due to outgassing from the sample, attributed to desorption of contaminants. The removal of contaminants can be seen in the spectra recorded at 75 ˚C, which displays more clear components with narrower peak widths when deconvoluted compared to spectra recorded at 25 ˚C. When the temperature is increased further, the C and O components remain the same, and not until a temperature of 400 ˚C do the C1s and O1s spectra change significantly. The C1s spectrum recorded at 400 ˚C primarily shows a broad peak associated with a C−C/C−H component, whereas the O1s spectrum displays a broad peak centered at 532.7 eV, mainly assigned to the underlying SiO2



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substrate.43 At this temperature the MLD film has greatly degraded. The degradation of the pMLD film by 400 ˚C is in agreement with previous reports on the structurally similar PMMA,44,45 for which a recent study showed that thermal annealing under high vacuum does not completely remove PMMA residues from graphene, but rather decomposes PMMA to amorphous carbon.46 The main component in the C1s spectrum (~284.9 eV) recorded herein at 400 ˚C could indeed originate from an sp2-rich amorphous carbon film. The atomic concentrations of C, O, and Si obtained from XPS plotted as a function of temperature are shown in Figure 6c. In line with the stability seen in the C and O high resolution spectra, the atomic concentration of O and C remains stable until at least 300 ˚C. At this temperature, the underlying SiO2 is for the first time detected (0.5%). At 350 ˚C the fraction of Si increases slightly to 2.2% with a concomitant decrease in the O content, despite O from SiO2 also being detected at this temperature. This suggests that C−O and/or C=O bonds are starting to degrade at this temperature. From the temperature dependent (C−H+C−C)/C=O ratio (assigned from the high resolution C1s spectra) shown on the secondary y-axis in Figure 6c, it is clear that this ratio starts to increase at temperatures > 300 ˚C, while the C=O/C−O/ C*−(C=O)−O ratio remains close to 1:1:1 up to 350 ˚C (data not shown). We propose that the ester bonds start to degrade around 350 ˚C, and when the temperature reaches 400 ˚C, the majority of these bonds have decomposed. At 400 ˚C, the main components detected are the underlying SiO2 substrate and carbon. The thermal stability of the films produced by pMLD exceeds what has been seen for polythiourea,19 polyurea,14 and crosslinked polyurea films,47 which were all also heat-treated under vacuum conditions. Polythiourea was stable to < 250 ˚C and polyurea to ~250 ˚C; while a crosslinked


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version of polyurea provided additional stability. That film was 90% degraded after vacuum annealing at 290 ˚C.

Figure 6. Thermal stability of EGM-DIP films analyzed in situ by XPS. In the high resolution (a) C1s and (b) O1s spectra recorded at 25, 75, 100, 150, 200, 250, 300, 350, and 400 ˚C, the dashed lines indicate the



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binding energy of the relevant components (see Figure 5). (c) Atomic concentrations of C, O, and Si, and the (C−H+C−C)/C=O ratio (open circles) plotted as a function of temperature. Traces of I detected up to 150 ˚C are not included in the figure.

Discussion From the wavelength-dependent analysis, the saturation curves, and the growth curves presented in Figures 1, 2, and 3, different insights can be gained into the mechanism of the pMLD process. First, controlled pMLD of EGM/DIP was successful only with an optical longpass filter with a cut-on wavelength of 250 nm. At lower wavelengths, significant photopolymerization of EGM takes place, and at higher wavelengths, above the photodissociation threshold of DIP, no deposition is observed. The n→σ* transition of the DIP C−I bond is observed at 256 nm (Figure S3), and thus matches the 250 nm band pass filter. The fact that both precursors are required to achieve a controlled MLD process suggests that EGM does not photopolymerize to a significant extent, and that homolytic cleavage of the C–I bond must be involved in the pMLD process. Second, in pMLD of EGM/DIP, saturation of EGM is achieved within 600 s, whereas the DIP step is slower and only saturates after 3400 s. This considerable difference may seem surprising, especially since the vapor pressure of DIP is higher than EGM: the pressure during a DIP pulse is usually ~140 mTorr compared to ~30mTorr for EGM. In a temperature-dependent control experiment, we found that the growth rate decreased by ~63% when increasing the temperature from 25 ˚C to 45 ˚C and by another 10% when going to 60 ˚C (Figure S6). The same trend has previously been observed for MLD of polyurea14 and polyamide.16 In the case of polyurea, the decrease in growth rate at 45 and 60 ˚C was very similar to what we see here. The behavior is explained by a precursor-mediated adsorption effect, where the residence time of the gas-phase precursor on the surface is important for efficient coupling reaction.16,48 By increasing the temperature, the lifetime of an adsorbed precursor state is shortened.


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The low growth rates suggest that precursor saturation of EGM and DIP is not achieved in pMLD at elevated temperatures, consistent with a precursor-mediated adsorption model. We propose that the difference in saturation time of EGM and DIP observed at room temperature is related to this physisorbed precursor state, combined with the recognition that C–I dissociation is central to the photoinitiated reaction chemistry. In other words, the reaction rate will depend on the surface concentration of C–I bonds. In (a) an EGM dose, the surface concentration of C–I bonds is given by the surface concentration of covalently-bound DIP molecules and is expected to be on the order of 1015 groups cm-2,49 whereas during (b) a DIP dose, the concentration of C–I bonds depends upon the concentration of the physisorbed precursor state of DIP. The pressure in the reactor during a DIP dose is below the vapor pressure of DIP, so we do not expect DIP to condense on the surface. Furthermore, in contrast to MLD of polyurea and polyamide, where hydrogen-bonding interactions can stabilize a precursor state between the reactant and the surface, for pMLD of EGM and DIP only the weaker van der Waals interactions are expected in the physisorbed state, which will reduce the lifetime of this precursor state. Hence, we may expect that the surface concentration of physisorbed DIP precursor in (b) will be lower than saturation coverage (< 1015 groups cm-2) and hence lower than covalently bound DIP in (a). This concentration difference likely accounts for the difference in the rate of saturation between DIP and EGM. Third, no significant deposition is observed for pMLD of either of the two precursors, EGM or DIP, alone, or for pMLD involving monofunctional MMA and bifunctional DIP as summarized in Scheme 2a, b, and c. These results are all consistent with a controlled MLD process.10 However, pMLD of bifunctional EGM combined with monofunctional IP (Scheme 2d) does exhibit an initial growth rate close to EGM/DIP, and with increasing number of cycles, the growth rate decreases and approaches a steady state growth rate. The occurrence of deposition using the monofunctional IP precursor was 21 ACS Paragon Plus Environment


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unexpected, since typically two functionalities are required for film growth in a step-growth polymerization. In contrast to the classic MLD polymers formed by step-growth polymerizations, pMLD propagates through a radical step-growth polymerization. The radicals are formed in the homolytic cleavage to the C–I bonds, creating I˙ and alkyl radicals. We propose the following growth mechanism to explain the film growth in pMLD of EGM and IP. In the photoactivated IP exposure, the C–I bond is homolytically cleaved, creating I˙ and n-propyl radicals. We suggest that the deposition is caused by the ability of both I˙ and n-propyl radicals to add to the ene precursor, EGM (Scheme 2d). If (i) I˙ addition occurs, the C–I bonds formed can undergo another photodissociation event in the following MLD half cycle and react with EGM enabling continued growth, whereas if (ii) n-propyl radical addition takes place, the chain propagation will terminate. Depending on the probability of the termination path, (ii), it is expected that the growth rate will level off and approach zero after a certain number of MLD cycles. The fact that the growth rate here approaches a steady state, non-zero value rather than zero is attributed to some degree of photopolymerization of EGM taking place with 250 nm UV-irradiation. Scheme 2. Proposed Mechanism of Iodo-Ene Photo-Activated Molecular Layer Deposition.


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O O

(a)

hv 250 nm Minor degree of photopoly- (b) merization

O

O

(c)

I

hv 250 nm

O

I

hv 250 nm I

I

No pMLD

No pMLD

O

O O (d)

O

I

I

hv 250 nm

O

I

(i)

Continued MLD

O

O

O

(ii)

O I

I hv 250 nm

I

Termination

I O O

O

I

O O

(e)

O

O

x

O

I

O

O x = 0 or 3

O O

O

O hv 250 nm (f)

I

I

I

I

I

I

I

H

To explore the extent of the termination pathway, we have derived Eq. 1, which models the film thickness (T) of EGM/IP as a function of the number of pMLD cycles (n). In the model, it is assumed that in each cycle, p is the fraction of C=C sites being terminated through (ii). We have additionally defined the term, GPCinit, which is the initial growth per cycle, and the term, GPCSS, which accounts for thickness increase from EGM units, added per cycle through photopolymerization. Eq. 1 is derived in the Supporting Information. ܶ = GPCୗୗ ቀ݊ −

ଵି௘ ష೛೙ ௣

ቁ + GPC୧୬୧୲

ଵି௘ ష೛೙ ௣

(Eq. 1)

Given the close resemblance between an EGM-DIP and an EGM-IP unit, GPCinit is here estimated to be equal to the measured GPC of EGM/DIP (3.7 ± 0.3 Å/cycle). Using Eq. 1, T is fitted to the experimental thicknesses of EGM/IP in Figure 3 using the least square method with p and GPCSS as variables. The optimized fitting gives p = 0.41 ± 0.05 and GPCSS = 0.81 ± 0.01 Å/cycle and the fit, as shown from the red, dashed curve in Figure 3, shows good agreement with the experimental data



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points. GPCSS has two contributions: (1) EGM units added through self-initiated photopolymerization and (2) EGM/IP MLD-growth from C=C groups introduced by EGM photopolymerization (see Scheme S3). A steady state growth rate of GPCSS = 0.81 Å/cycle is not unreasonable taking into account that a self-limiting growth (terminated growth after a few cycles) of ~0.8 Å was observed for pMLD of EGM alone, and that the additional EGM/IP growth from new C=C surface sites is also included in GPCSS. The value of p = 0.41 ± 0.05 that arises from the model suggests that ~40% of the growing polymer chains terminate every cycle, and that the remaining ~60% are capped by I. The lower probability of the chain termination step suggests that attachment of n-propyl radical to the alkene group in EGM is less facile than attachment of an I˙. We hypothesize that the difference is related to instability of the n-propyl radical. The n-propyl radical readily undergoes fragmentation to form propene, ethylene, and methyl radical (Scheme S4).50,51 Additionally, some of the as-formed npropyl radicals are known to be in a higher vibrational excited state, and these energized propyl radicals are prone to abstract hydrogen and form propane.26 These secondary photodissociation reaction products lower the concentration of n-propyl radicals and may explain the 41:59 reaction ratio between addition of the n-propyl radical and I˙. The two addition pathways, n-propyl radical and I˙ addition, explain the growth behavior seen for pMLD of EGM/IP. However, it also suggests that the mechanism initially proposed for pMLD of EGM/DIP is not as simple as that presented in Scheme 1. We therefore revisit the mechanism of pMLD of EGM/DIP considering the mechanistic aspects learned from pMLD of EGM/IP (Scheme 2d and e). During a DIP half cycle, both I˙ and 1-iodopropyl radical can add to surface bound EGM and in both cases give an iodo-capped surface site. In the subsequent half cycle, EGM can react with either of these species to continue the MLD process. Notably, controlled MLD is expected irrespective of which of these two pathways occurs. The only difference is the chain length of the alkyl fragment in the polymer 24 ACS Paragon Plus Environment

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backbone, as shown in the product of Scheme 2e. Assuming the presence of these two operating pathways, from the XPS results presented in Table 1 and Figure 5, both the C/O and (C−H+C–C)/C=O ratios may independently be used to estimate of the average alkyl chain length, x, in the polymer backbone. Here x = 0 in the case of I˙ addition, and x = 3 for addition of 1-iodopropyl radical. From the C/O ratio we derive a value of x = 0.96, whereas the (C–H+C–C)/C=O ratio gives x = 1.0. Both values are very close to x = 1. This corresponds to 1-iodopropyl units being added in ~33% of the repeating units. Interestingly, these percentages are in good agreement with 41% n-propyl radical addition (p = 0.41) found from modelling the EGM/IP process, and they support our hypothesis that the pMLD process of EGM/DIP operates through two parallel pathways. In fact, a minor degree of EGM photopolymerization would appear as the I˙ addition pathway, meaning that in reality the 33% is presumable slightly higher and closer to the 41% obtained from the modeling. A final mechanistic aspect which should also be addressed is the instability of the 1-iodopropyl radical mentioned previously. Studies of the photolysis of 1-iodopropane at 265 nm have shown that the propyl radical formed upon homolytic cleavage of 1-iodopropane undergoes fragmentation to give the secondary reaction products ethylene, methyl radical, and propylene (Scheme S4).51 Herein, it may be expected that 1-iodopropyl radical can fragment the same way to give ethylene, iodomethyl radical, and 3-iodoprop-1-ene (Scheme 2f). Iodomethyl radical may add to EGM and participate in the MLD cycle analogous to the 1-iodopropyl radical of Scheme 2e. Let us assume that this is a third viable pathway for pMLD of EGM/IP and EGM/DIP. In the EGM/IP model presented in Eq. 1, termination events through addition of the n-propyl radical or the fragmented methyl radical cannot be distinguished. Therefore the model implies that 59% of the reactions proceed through addition of I˙, whereas the remaining 41% of the reactions are termination



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events through addition of the n-propyl radical or the fragmented methyl radical. If we assume that 59% of pMLD of EGM/DIP also proceeds through addition of I˙ while 41% of the reactions occur through addition of the 1-iodopropyl radical or the fragmented iodomethyl radical, we will have three contributions to the value of x (average alkyl chain length): 59% is addition of I˙ (x = 0) and 41% includes addition of 1-iodopropyl radicals (x = 3) and iodomethyl radicals (x = 1). From the XPS results presented earlier we found x = 0.96 based on the C/O ratio, which corresponds to 28% 1iodopropyl radical and 13% iodomethyl radical, whereas the (C–H+C–C)/C=O ratio gave x = 1.0 corresponding to 30% 1-iodopropyl radical and 11% iodomethyl radical. Based on the discussion given above, we have found that the mechanism behind pMLD of EGM/DIP primarily has contributions from two different pathways as outlined in Scheme 2e: (1) addition of I˙ and (2) addition of 1-iodopropyl radical. The third possible pathway involving the fragmented iodomethyl radical has a minor contribution of 11–13% to the overall pMLD mechanism of EGM/DIP. We emphasize that this new class of MLD polymerizations, i.e. pMLD, takes advantage of radical reactions to overcome the otherwise sluggish kinetics of gas phase coupling reactions of organic molecules. Due to the radical species involved in pMLD, it is challenging (if not impossible) to achieve full control and to completely avoid side reactions. However, we are able to achieve a controlled deposition of acrylate polymers, with saturation behavior for each of the two precursors and a constant growth rate. Moving forward, we are interested in studying the effect of more stabilized alkyl radicals formed upon photo-induced homolysis of C–I bonds. Furthermore, we also aim to explore the chemical versatility of pMLD. pMLD may in fact lead the way for the realization of a new diverse class of nanoscale organic films not attainable through traditional MLD coupling chemistries. We expect that the unique molecular control achieved by MLD will find renewed interest and applications with a larger tool-box of MLD coupling chemistries available. 26 ACS Paragon Plus Environment

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Conclusions In this work, we demonstrate a new type of MLD chemistry, that of photoactivated radical stepgrowth polymerization. Successful pMLD is carried out using EGM and DIP sequentially exposed to a UV-irradiated silicon substrate. The pMLD process occurs by constant growth rate, and both monomers show self-limiting saturation behavior. The organic surface-tethered films exhibit excellent thermal stability and withstand solvent stability testing. Characterization by FTIR spectroscopy shows strong absorption from the ester bonds in the polymer backbone, while full conversion of C=C in EGM monomers is confirmed by the absence of a C=C stretch mode. XPS analysis confirms that the polymer films contain the expected atomic and chemical components, but surprisingly the C/O and (C–H+C– C)/C=O ratios suggest that the mechanism behind pMLD of EGM and DIP is more complex than a simple step-growth polymerization. Interestingly, pMLD of EGM and monofunctional IP does not result in immediate polymer termination, which suggests that I˙, obtained from the photo-induced homolytic cleavage of DIP, may also add to EGM monomers and participate in the pMLD process. We propose that pMLD of EGM and DIP also occurs through a combination I˙ and 1-iodopropyl radical addition. Modelling the growth of EGM/IP suggests that 59% reaction proceeds through I˙ addition, which is in good agreement with XPS findings of the EGM/DIP process.



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Associated Content Supporting Information: MLD deposition conditions, nitrogen purge optimization study, AFM surface roughness analysis, UV-vis of DIP, XRD analysis, computational methods for DFT calculations, DFT calculated IR spectra of EGM and EGM-DIP and the experimental spectrum of EGM monomer, temperature-dependent growth of EGM/DIP pMLD, and modeling pMLD of EGM/IP. Acknowledgments This work was supported by the National Science Foundation (CHE-1607339). M.L. acknowledges support from the Villum Foundation and the Carlsberg Foundation, D.S.B. an NSF Graduate Research Fellowship, and T.S. a fellowship from Becas Chile Conicyt. Part of this work was performed at the Stanford Nano Shared Facilities (SNSF).

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