Effect of Backbone Chemistry on the Structure of Polyurea Films

Subscriber access provided by UNIV OF CALIFORNIA SAN DIEGO LIBRARIES

Article

Effect of Backbone Chemistry on the Structure of Polyurea Films Deposited by Molecular Layer Deposition David S. Bergsman, Richard G. Closser, Christopher J. Tassone, Bruce M. Clemens, Dennis Nordlund, and Stacey F. Bent Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b04530 • Publication Date (Web): 01 Jan 2017 Downloaded from http://pubs.acs.org on January 4, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

Chemistry of Materials is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 17

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

Chemistry of Materials

Effect of Backbone Chemistry on the Structure of Polyurea Films Deposited by Molecular Layer Deposition David S. Bergsman†, Richard G. Closser‡, Christopher J. Tassoneǁ, Bruce M. Clemens#, Dennis Nordlundǁ, and Stacey F. Bent†‡#* †

‡

#

Department of Chemical Engineering, Department of Chemistry, Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, United States. ǁSLAC National Accelerator Laboratory, Menlo Park, California 94025, United States. ABSTRACT: An experimental investigation into the growth of polyurea films by molecular layer deposition was performed by examining trends in the growth rate, crystallinity, and orientation of chains as a function of backbone flexibility. Growth curves obtained for films containing backbones of aliphatic and phenyl groups indicate that an increase in backbone flexibility leads to a reduction in growth rate from 4 Å/cycle to 1 Å/cycle. Crystallinity measurements collected using grazing incidence x-ray diffraction and Fourier transform infrared spectroscopy suggest that some chains form paracrystalline, out-of-plane stacks of polymer segments, with packing distances ranging from 4.4 Å to 3.7 Å depending on the monomer size. Diffraction intensity is largely a function of the homogeneity of the backbone. Near-edge x-ray absorption fine structure measurements for thin and thick samples show an average chain orientation of ~25° relative to the substrate across all samples, suggesting that changes in growth rate are not caused by differences in chain angle, but instead may be caused by differences in the frequency of chain terminations. These results suggest a model of molecular layer depositionbased chain growth in which films consist of a mixture of upward growing chains and horizontally aligned layers of paracrystalline polymer segments.

Introduction In recent years, the use of molecular layer deposition (MLD) has seen a surge of interest for applications that require ultrathin organic and hybrid organic-inorganic materials, such as in photolithography1–4, lithium ion batteries5–7, and catalysis8,9. Similar to its purely inorganic counterpart atomic layer deposition (ALD), MLD is a vapor phase technique that uses an alternating sequence of self-limiting reactions of precursors to grow a film on a substrate. The self-limiting nature of each precursor reaction allows for the deposition of highly conformal films on even high aspect ratio surfaces, with precise control over the film thickness. When organic monomers are used as precursors, the backbone of each monomer and the linkage chemistry between monomers can also be independently selected from the vast library of organic functionalities, allowing for the precise tuning of film properties10. In this way, MLD is ideally suited for the creation of ultrathin organic layers, and has been used to deposit a variety of

different types of polymers, such as polyimide11,12, polyurea13,14, polyamide15,16, polyurethane17, 18 19 polythiourea , polyazomethine , polyester20, and polyethylene-dioxythiophene21. In addition, monomers can be used in conjunction with ALD precursors to create hybrid films, which combine the chemical flexibility of organic layers with the structural and electrical properties of inorganic layers to create unique materials22. MLD is distinctive among liquid and vapor phase polymerization techniques in that each layer of monomers has the opportunity to self-assemble and self-arrange before the next layer is deposited. Indeed, recent work has shown that monomers and small molecules near a free surface can adopt unique configurations compared to the bulk, due to the increased freedom of motion near the polymerair interface23,24. However, it is not obvious what structure the MLD films would be expected to form. There are a number of different factors that contribute to the resulting polymer structure, such as the molecular weight of the final polymer chain,

1 ACS Paragon Plus Environment


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

the packing density of chains on the substrate surface, the excluded volume of each monomer, the inherent flexibility of each monomer, the deposition temperature, and the strength of intermolecular interactions such as van der Waals forces or hydrogen bonding. Given the wide range of interactions involved in this layer-by-layer process, films could adopt any number of different structures, ranging from tightly packed SAM-like layers to multiconformational organic films. The ability to control these structures can have a tremendous impact on many properties of the film, such as the thickness, film density, crystallinity, and thermal or electrical conductivity. The morphology of these organic layers can also impact the properties of hybrid materials, and previous work has attempted to take advantage of different MLD chemistries in order to prevent cross-linking3, decrease thermal conductivity25, and increase the density of carbon within films26. However, despite recent efforts to study and control the morphology of organic and hybrid organic-inorganic films deposited by MLD27,28, a greater understanding of the origin of these properties is still needed to enable the use of this technique in many nanotechnology applications. In this work, we examine, in detail, the growth behavior and structure of polyurea films deposited by molecular layer deposition, both during their initial surface nucleation and in the bulk. Polyurea MLD films of various backbone chemistries were studied with a combination of ex situ variable angle spectroscopic ellipsometry (VASE), Fourier transform infrared (FTIR) spectroscopy, grazing incidence x-ray diffraction (GIXD) and angledependent near edge x-ray absorption fine structure (NEXAFS). Trends in the film properties as a function of backbone flexibility were then used to gain insight into the film morphology. Previous work27–31 has shown that the choice of backbone chemistry can affect film properties. Here, we provide a comprehensive study of the effect of backbone flexibility on the structure of films grown by MLD, using polyurea films as a model. We demonstrate that both intermolecular forces between chains and monomer backbone chemistry have a significant impact on the film growth behavior, orientation, and structure. We further provide additional evidence that the use of more flexible monomers leads to an increase in the number of chain terminations. Insight into these properties will allow for improved design of new MLD chemistries and better control over the properties of films deposited by MLD. Materials and Methods

Page 2 of 17

1,4-phenylene diisocyanate (PDIC), 1,4diisocyanato butane (BDIC), 1,2-ethylenediamine (ED), 1,4-diaminobutane (BD), and pphenylenediamine (PD) were purchased from Sigma Aldrich and used as received. All reagents were stored in a nitrogen-purged glove box until use. Films were deposited on silicon (100) wafers purchased from WRS Materials with a 1.6 nm native oxide, as measured by VASE. Before use, silicon wafers were first cleaned for 15 minutes in a piranha solution with a 7:3 ratio of 98% sulfuric acid/30% hydrogen peroxide before being rinsed and stored in deionized water until use (Caution: fresh piranha solution is hot and extremely corrosive. Proper training and extreme care should be taken when preparing and handling piranha solution.). Immediately before film deposition, all samples were dried with compressed air and then cleaned for 15 minutes in a Novascan PSD Series Digital UV Ozone System to remove any remaining organic contaminants. Samples were then introduced to the reaction chamber for deposition. MLD films were deposited as described previously14, using one of two hot-wall flow reactors of similar design, both pumped by a Leybold Trivac rotary vane pumps with a base pressures near 5 mTorr. The bodies of the reaction chambers were kept at room temperature (RT) in order to maximize their growth rates14 and promote increased ordering32, but some individual monomers were heated in order to achieve sufficient vapor pressures for deposition (PDIC, BDIC, ED, BD, and PD were heated to 35 °C, RT, RT, 27 °C, and 45 °C respectively). All chemistries used the same urea linkage, which has been explored previously14, in order to impart consistency across the films. Details of the deposition conditions are included in the supporting information. Immediately after deposition, VASE was performed using a J. A. Woolam Co. alpha-SE spectroscopic ellipsometer with a spectral range of 300 to 900 nm. Measurements were taken at incidence angles of 65°, 70°, and 75° relative to the surface normal. Thickness was modeled by employing a Cauchy model with a constant refractive index of 1.8 and zero extinction term for all of chemistries. Though this refractive index is unphysical, the model and refractive index were chosen because they accurately modeled film thickness, as verified by atomic force microscopy scratch tests (see Supplemental Table 1) and X-ray reflectivity measurements (Supplemental Table 2). Uncertainty in film thickness from ellipsometry is estimated to be around ± 0.2 Å, based on repeat


Page 3 of 17

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


scans of the same sample over several days. Each data point on the growth curves represents an independent run. FTIR spectroscopy measurements were performed using a Nicolet iS50 FTIR spectrometer with a diamond attenuated total reflectance (ATR) plate. Spectra were taken with 200 scans at 2 cm-1 resolution with equivalent UV-ozone treated Si substrates used as a background. Grazing incidence and symmetric XRD measurements were performed on Beamline 11-3 at the Stanford Synchrotron Radiation Lightsource (SSRL) in a helium-purged chamber with Kapton windows, using a MAR345 imaging plate area detector with x-ray energy of 12.7 keV. The detector distance of around 400 mm and detector orientation were calibrated using a LaB6 standard and confirmed by comparison to the Si(111) at 2.0 Å-1 present in several samples. Grazing incidence scans were taken at an incidence angle of 0.12° with respect to the x-ray beam, while symmetric scans were performed at an angle corresponding to the peak of interest, as described previously33. Uncertainty in d-spacings was estimated to be around 0.1 Å based on taking multiple scans of the PDIC+ED chemistry and extracting the root mean squared error in the d-spacing for that peak. Angle-dependent NEXAFS was performed at SSRL on beamline 8-2, in a chamber pumped to below 10-8 Torr and held at room temperature. Measurements were collected at a grazing incidence angle of 20°, 35°, 55°, 70°, and 90° relative to the substrate surface at three different spots on each sample (following the geometrical definitions of Stöhr and Outka34). Each spectrum was then fit to a series of asymmetric Gaussians and exponentially decaying step functions, as shown in Supplemental Tables 3-12, based on previous fits of polyurea NEXAFS spectra35. The peak areas were then fit using a least squares minimization to the threefold or higher symmetry equation for orbital orientation described by Stöhr and Outka34, in order to determine bond orientation. Uncertainty in bond orientation is estimated to be around 3°, based on a sensitivity study of the fitted intensities when varying the peak positions and widths. Full details for spectra correction and fitting are included in the supplemental information.

flexibility36 (Table 1) were used to deposit polyurea MLD films on silicon substrates, and their thicknesses were measured with VASE as shown in Figure 1. Monomers were chosen based on the following criteria: 1) the resulting chemistries spanned a range of different flexibilities, 2) the individual monomers were relevant to and used in existing and current organic and hybrid organicinorganic chemistries of interest [2,22,37], 3) the monomers were small enough to have sufficient vapor pressure to achieve growth, and 4) the monomers could be dosed into a room temperature chamber without physisorbing and crystallizing on the substrate. Table 1: Depiction of monomers used and the resulting polymer chemistry. Monomer A

p-phenylene diisocyanate (PDIC)

1,4diisocyanato butane (BDIC)

Monomer B

Polymer

p-phenylene diamine (PD)

(PDIC+PD)

1,2-ethylene diamine (ED)

(PDIC+ED)

1,4-diamino butane (BD)

(PDIC+BD)

1,2-ethylene diamine (ED)

(BDIC+ED)

1,4-diamino butane (BD)

(BDIC+BD)

Polymers are organized by increasing backbone flexibility moving from the top to the bottom of the table.

Results Insight into the growth behavior of polyurea MLD films was obtained by examining the effect of the monomer backbone on the packing and orientation of the resulting chains. To that end, diisocyanate and diamine monomers with increasing backbone



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

Page 4 of 17

After reaching steady state, all of the studied chemistries experienced a constant growth rate for up to 32 cycles. However, the growth rates varied widely between the most flexible and most rigid backbones. MLD films with the most rigid monomers experienced growth per cycle (GPC) rates around 4 Å/cycle, which agrees well with the growth rate previously recorded for the mixed backbone and phenyl-phenyl backbone polyureas14,27. In contrast, the chemistry with the most flexible backbone yielded a relatively low growth rate of 1 Å/cycle. All of the MLD chemistries also showed surface enhanced growth during the first few cycles of deposition, with growth rates approaching 6 – 8 Å/cycle.

Figure 1: Growth curves for the MLD chemistries a) PDIC+PD, b) PDIC+ED, c) PDIC+BD, d) BDIC+ED, e) BDIC+BD. Each MLD process shows linear growth per cycle rates of 3.9 Å/cycle, 4.2 Å/cycle, 3.3 Å/cycle, 2.4 Å/cycle, and 1.0 Å/cycle respectively. Surface-enhanced growth regions are shown as dashed lined.

Figure 2: FTIR spectroscopy data taken for MLD films of a) 32x[PDIC+PD], b) 24x[PDIC+ED], c) 32x[PDIC+BD], d) 32x[BDIC+ED], and e) 32x[BDIC+BD]. Peak assignments are summarized with bracket annotations. The inset (data shown in same top-to-bottom order) emphasizes the red-shift and blue-shift of the v(C=O) and δ(N-H) peaks respectively that results from the varying hydrogen bond lengths of the most rigid to the most flexible polymers.


Page 5 of 17

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


The successful reaction between each pair of monomers was confirmed by FTIR spectroscopy (Figure 2). Each of the spectra shows peaks characteristic of the urea linkage, including a C=O stretch at 1600-1700 cm-1, an N-H urea wag around 1500-1600 cm-1, and an N-C-N stretch at 1300-1305 cm-1. Assignment of the remaining peaks, including a broad N-H stretch around 3300 cm-1, C-H stretch features around 2800-3000 cm-1, an NH2 scissor mode near 1605 cm-1, a sharp ring stretching mode around 1500 cm-1, a CH2 bend around 1450 cm-1, and an aromatic ring mode at 1404 cm-1, were based on previous reports of the polyurea MLD films13,14. These previous reports suggested that the sharp peak around 1500 cm-1 in the more rigid chemistries is the N-H wag of urea, while the broader peak around 1500-1600 cm-1 is the ring stretching mode. However, we have reversed these assignments based on the more extensive data collected here for the following reasons: 1) the aliphatic chemistries do not show the sharp peak around 1500 cm-1, and 2) the broad peak shows a much more significant blue shifting as a result of hydrogen bonding than the sharp peak, as discussed below. Closer examination of the C=O carbonyl stretch and N-H urea wag shows an overall trend of increased red-shifting and blue-shifting of these peaks, respectively, with increasing chain flexibility (inset to Figure 2). These shifts have been associated with the degree of hydrogen bonding between urea groups38,39 with a smaller splitting between the carbonyl and N-H wag being associated with stronger hydrogen bonding in the film. Hence the most flexible chain, which according to curve (e) has the smallest splitting, yields the strongest hydrogen bonding. Previous work40 has also shown that the carbonyl stretch of urea groups absorbs at different frequencies depending on their degree of crystallinity, with non-hydrogen bonded urea groups, disordered hydrogen bonded urea groups, and ordered hydrogen bonded urea groups absorbing at roughly 1690 cm-1, 1650 cm-1, and 1630 cm-1, respectively. Based on peak shapes and positions, this would indicate that the chemistries with backbones containing both rigid and flexible segments (PDIC+ED and PDIC+BD) experience a mix of free and disordered hydrogen bonded urea groups, while PDIC+PD experiences primarily disordered hydrogen bonding and the all-aliphatic backbone chemistries (BDIC+ED and BDIC+BD) experience primarily ordered hydrogen bonding. GIXD was performed on films grown with each of the chemistries. Representative diffraction patterns are shown in Figure 3(a-e). Here, q is the scattering vector, with qz representing the purely out-of-plane

component of q and qxy representing the in-plane x and y components of q, flattened into one image. The variable χ will represent the azimuthal angle, defined to be 90° when q is purely out-of-plane. All samples, including a clean silicon wafer (not shown), showed the presence of a diffuse peak at q = 2.0 Å-1, χ = 33° assigned to Si(111)41, with some samples also showing isotropic silicon powder diffraction rings. Notably, each of the MLD films also shows a diffuse peak, oriented roughly out-ofplane (in the qz direction with little qxy component). This peak is unchanged upon in-plane φ rotation of the sample, suggesting fiber symmetry.

Figure 3: GIXD patterns of films consisting of a) 48x[PDIC+PD], b) 45x[PDIC+ED], c) 50x[PDIC+BD], d) 48x[BDIC+ED], e) 24x[BDIC+BD], with film thicknesses of 280.1 Å, 288.43 Å, 65.0 Å, 155.0 Å, and 88.8 Å, respectively. Color bar scaling has been adjusted independently for each plot to show contrast between the relevant peak and the background. Annotations (white) outlining the approximate q-value of silicon powder diffraction and the location of the out-of-plane MLD peak are shown in (a). f) Integrated intensity of the out-of-plane diffraction peaks shown in a-e, labeled as i-v respectively. The integrated region is outlined (red) in (a). Intensity of the BDIC+BD 2θ scan was scaled down by a factor of 10.



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

From these patterns, the data was integrated with respect to q over a pie segment of the azimuthal angle χ from 80° to 100° and normalized with respect to film thickness, in order to determine the peak positions for the purely out-of-plane qz direction (Figure 3f), and the resulting intensities were fit to Lorentzian peaks in order to determine their peak positions, areas, and full width half maxima in roughly the qz direction (Table 2). Since the purely qz diffraction direction is unobservable in GIXD33 and the peak intensity is too low to easily see in a symmetric geometry, a combination of grazing incidence and symmetric geometries were performed to create a rocking curve in χ for the outof-plane peaks (Figure 4)33, highlighting the degree of out-of-plane texture. Only three of the chemistries (PDIC+PD, BDIC+ED, and BDIC+BD) had sufficient out-of-plane intensity to be observable in the symmetric configuration and could be used to make the combined rocking curves. These rocking curves show that the peaks are indeed centered out-of-plane. Moreover, the peak widths in the χ direction (48.6°, 34.2°, and 19.3° for PDIC+PD, BDIC+ED, and BDIC+BD respectively) decrease with increasing flexibility, suggesting greater out-of-plane alignment for the most flexible chemistries. Table 2: Out-of-plane peak fits for each of the MLD chemistries. q

d

Area

FWHM

-1

4

4.4 Å

8.4 x10

0.2 Å

-1

4

0.5 Å

-1

PDIC+PD

1.4 Å

PDIC+ED

1.5 Å

-1

4.2 Å

4.2 x10

PDIC+BD

1.5 Å

-1

4.1 Å

20.1 x10

0.6 Å

BDIC+ED

1.6 Å

-1

3.9 Å

7.0 x10

4

0.2 Å

-1

BDIC+BD

1.7 Å

-1

3.7 Å

81.5 x10

0.2 Å

-1

4

4

-1

Includes q-values, d-spacings, peak areas, and full width half maxima associated with the out-of-plane peak for each of the MLD chemistries.

Of the five chemistries studied, the three chemistries with homogeneous backbones (PDIC+PD, BDIC+ED, BDIC+BD) show a much lower full width half maximum for the out-of-plane peak than the two chemistries with mixed backbones (PDIC+ED and PDIC+BD) (as evident in Figure 3(f)), suggesting that they are better able to form well-ordered out-of-plane structures. This agrees well with polymer literature: having a mix of backbone chemistries makes it more difficult for the similar-backboned regions to align and form wellpacked domains42. The C=O stretch peak position in the FTIR spectra also reflects this trend: the

Page 6 of 17

homogenous backbone polymers showed more ordered hydrogen bonding than did the mixed backbone polymers. The observation of a diffuse out-of-plane diffraction peak in fiber symmetry, with a d-spacing around 4 Å and a full width half maximum ranging from 0.2 to 0.5 Å-1, suggests that at least some of the chains are forming a paracrystalline out-of-plane stack of polymer segments43, without any strong, ordered lateral interactions. The measured dspacing represents the distance between the packed chains, which provides rationale for the trend in reduced d-spacing with increased chain flexibility: the bulkier, more rigid backbones take up more volume and prevent the chains from packing as tightly as the thinner, aliphatic-backboned chains. These d-spacings are also roughly consistent with literature values for the d-spacings of each chemistries' respective monomers44–48. The origin of this order and its implication on the growth behavior of films will be discussed in more detail in the discussion section. Among the systems studied, the BDIC+BD chemistry is unique in that it shows additional diffraction features (Figure 3e), with an out-of-plane peak at q=0.8Å-1, two diffuse peaks around 1.1Å-1 and 1.2Å-1, χ = 50°, and an amorphous ring at q=1.64Å-1, with extra intensity at χ = 33° (Figure 3e). This richer GIXD pattern is the result of the formation of a more-complete 2D lattice, in contrast to the 1D lattice seen in the other chemistries. The fact that several of the peaks extend over multiple χ directions suggests that this chemistry also forms crystallites aligned in directions other than the outof-plane direction, likely due to its high degree of flexibility and ability to crystallize. The out-of-plane peak around q=0.8Å-1 is suspected to be a superlattice peak resulting from a slight mismatch between chain layers, since its d-spacing is roughly twice as large as the prominent peak discussed above. Literature values for the peaks of diureidobutane49 are also similar to those observed for the BDIC+BD chemistry, which is similar in relative composition; however, a complete crystallographic analysis of the BDIC+BD chemistry was not performed in this study.


Page 7 of 17

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


measurements. Relatively thick (24-48 cycles) and thin (4-8 cycles) films of each chemistry were grown and taken to SSRL for ex situ analysis. Investigation of thicker samples allows for a better characterization of bulk chain properties, while study of thinner samples allows examination of films during the first few cycles of growth. Representative normalized spectra for the carbon Kedge and nitrogen K-edge excitations of PDIC+PD are shown in Figure 5. The complete sets of spectra for all of the chemistries studied are presented in the supplemental information.

Figure 4: Combination of symmetric and grazing incidence rocking curves along χ for a) PDIC+PD, b) BDIC+ED, and c) BDIC+BD films, based on Baker et 33 al. The insets show the grazing incidence (left) and symmetric geometry (right) diffraction patterns. Peaks exhibit full width half maxima, of a) 48.6°, b) 34.2°, and c) 19.3°.

Each of the films showed the expected resonances for their associated chemistry when compared to literature values35. The following text briefly highlights these resonances, using the notation C1s→1π*C=O, where "C1s" represents the origin of the electron (a 1s shell electron from carbon and "1π*C=O" represents the transition orbital (the antibonding 1π* orbital of the C=O bond). In the carbon K-edge spectra, clear pre-step edge resonances include the C1s→1π*C=C transitions around 286 eV and 287 eV for chemistries containing phenyl backbones and a C1s→1π*C=O around 290 eV for all chemistries investigated. Other peaks, previously assigned to C1s→2π*C=C and C1s→3s resonances35 are observable as shoulders to the C1s→1π*C=O peak, but were not rigorously fit in intensity calculations. Post-edge C1s spectra included the C1s→1σ*C=C resonance around 294 eV and 303 eV, the C1s→1σ*C-N resonance around 300 eV, and the C1s→1σ*C=O resonance around 304 eV. Previous work also assigns a C1s→1σ*C=C transition to the same energy as the C1s→1σ*C=O resonance35. The N1s spectra include the N1s→1π*N-C resonance around 403 eV and the N1s→1σ*N-C resonances around 408 eV and 412 eV, though minor shoulder transitions of N1s→3s, N1s→3p, and N1s→4p are also observed. Surprisingly, the BDIC+ED chemistry showed traces of phenyl group character, likely a result of minor cross-contamination between different chemistries during deposition. This was verified though closer inspection of the FTIR spectroscopy data, which also shows a trace ring stretch near 1500 cm-1 for the BDIC+ED chemistry. This relatively minor cross-contamination is not expected to affect the average orientation significantly.

The average orientation of the chains was examined using angle-dependent NEXAFS



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

Page 8 of 17

though the C1s→1π*C=C transition to see how the packing of chains may affect their alignment. The resulting bond orientations extracted from the NEXAFS data and analysis are given in Table 3.

Figure 5: Plots of the carbon K-edge and nitrogen Kedge NEXAFS spectra of a bulk (24 cycles) film of PDIC+PD taken at different incidence angles.

Scans taken at different incidence angles show that spectroscopic transitions associated with several constituent bonds—that of the phenyl rings in the aromatic backbones, the C=O π bond, the CN σ bond, and the C=O σ bond—exhibit preferential orientation within the film. Based on the work of Stöhr and Outka34, the changes in peak intensity of a particular bond as a function of incidence angle were used to determine the orientation of the bond with respect to the surface normal. Since the urea group is aligned to be oriented with respect to the polymer backbone, the bonds associated with the urea moiety can be used to determine the orientation of the chains. Due to the inherent overlap of the second C1s→1σ*C=C with the C1s→1σ*C=O in the C 1s spectra, the orientation of the C=O σ bonds could not be analyzed. The orientation of the phenyl rings was also examined

Based on the angle-dependent NEXAFS spectra, the bonds within each of the chemistries studied show only mild preferential orientation, having average bond orientations within roughly 10° of the magic angle (azimuthally averaged bonds with an average orientation of 54.7° relative to surface normal will show no variation in intensity with xray incidence angle)34. This result is consistent with previous NEXAFS studies performed on polyamide films, which showed similar intensity changes and bond orientations50. In general, the orientation of the C-N σ bonds of the urea group have a slight preference for a horizontal alignment, with bond orientations around 65° relative to the surface normal, while the π bonds of the C=O and N-C have a slight preference for a vertical alignment (around 45°). The phenyl rings show slight preference for lying flat, with a π bond orientation near 50°. Comparison between the carbon K-edge and nitrogen K-edge spectra show very good overall agreement, with C1s→1σ*C-N and N1s→1σ*N-C, as well as the C1s→1π*C=O and N1s→1π*N-C bond orientations within 5° (± 7°) and 1° (± 2°) respectively. The greater mismatch between the σ* bonds than between the π bonds is understandable, since π peaks are separate from the step edge and therefore much easier to fit. This is particularly true for the BDIC+ED and BDIC+BD thin films, which show relatively flat C 1s absorption features immediately after the step edge. Interestingly, there is not a significant trend in the bond orientations for the different chemistries. Indeed, the C1s→1σ*C-N and N1s→1σ*N-C resonances imply that all chemistries grow at an average angle of approximately 25° relative to the surface parallel (90° minus the average of the C-N bond orientations). This is a surprising result, considering the vastly different growth rates of the various MLD chemistries. There is also only a minor difference in the average bond orientations between the thick and thin films (3° ± 3°), in spite of the increased growth rate near the surface observed by VASE.


Page 9 of 17

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


Table 3: Bond orientations relative to the surface normal, based on angle-dependent NEXAFS fits. Thick Films (24-48 cycles) C-N → σ*C-N

C=O → π*C=O

N-C → π*N-C

C-C → π*C=C

C 1s

N 1s

C 1s

N 1s

C 1s

24x[PDIC+PD]

64°

64°

48°

44°

51°

45x[PDIC+ED]

61°

66°

50°

47°

48°

32x[PDIC+BD]

61°

60°

51°

51°

52°

48x[BDIC+ED]

65°

61°

47°

46°

–

32x[BDIC+BD]

65°

62°

43°

45°

–

C=O → π*C=O

N-C → π*N-C

C-C → π*C=C

Thin Films (4-8 cycles) C-N → σ*C-N C 1s

N 1s

C 1s

N 1s

C 1s

4x[PDIC+PD]

67°

65°

47°

45°

49°

4x[PDIC+ED]

68°

59°

52°

51°

49°

4x[PDIC+BD]

61°

60°

51°

51°

50°

8x[BDIC+ED]

78°

62°

45°

45°

–

4x[BDIC+BD]

79°

62°

42°

42°

–

Discussion There are several trends that are directly evident from the VASE, FTIR, GIXD, and NEXAFS results presented above. First, as the backbone flexibility of a particular chemistry increases, the resulting growth rate of that chemistry decreases. This trend has been observed previously in literature, both from a comparison between very rigid and very flexible backbones37 and as a function of aliphatic backbone length28 in the hybrid alucone MLD system. It is worth emphasizing that it is chain flexibility and not extended chain length that follows the trend observed here. The fully-extended repeat unit length for the five chemistries is estimated to be roughly 15.2 Å, 12.6 Å, 16.5 Å, 14.6 Å, and 17.0 Å in order of increasing flexibility (PDIC+PD to BDIC+BD), based on previous DFT calculations14. Since the growth rate of the chemistries follows a trend in backbone flexibility and not in backbone length, then the correlation must be attributed to backbone flexibility. Second, all chemistries exhibit some out-of-plane paracrystalline ordering in fiber symmetry and follow a trend of reduced d-spacing with increased chain flexibility. Third, those chemistries with homogeneous backbones experience a greater degree of crystallinity than those with a mix of aliphatic and phenyl groups. Fourth, all chemistries display roughly the same average bond orientation,

both in their first few cycles of growth and in the bulk, which suggests that the average angle of growth for the chemistries relative to the substrate (the "growth angle") is also all the same. We highlight that NEXAFS is a technique in which each atom contributes equally to the total intensity and the resulting trend is a linear average, thus we cannot easily distinguish between chains that are perfectly aligned at a particular angle and an ensemble of chains with orientations centered near that angle. The four trends identified in this work are highlighted in Figure 6.



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

Page 10 of 17

evidence has been presented to confirm its occurrence in the absence of other effects.

Figure 6: Summary of the trends in film properties as a function of backbone flexibility.

The differences in the observed growth rates of the chemistries studied here are of particular interest, since many factors have been proposed to affect MLD growth rates10. In general, there are two categories of effects that are thought to alter the growth rate of a particular chemistry: the amount of material being deposited per cycle (affected by the ability of the process to reach saturation and the surface site density) and the resulting structure of that material after being deposited (its density and the orientation of the chains relative to the surface). In the following, these two categories are discussed in more detail. For the first category, i.e. the amount of material deposited, the following factors play a role: 1) saturation and 2) surface site density. Saturation, or the ability to react with every available reactive site in each cycle, is often verified during the process characterization in order to mitigate the effect of subsaturation on the growth rate, as was done in the current study. There may be some reactive sites that remain unreacted due to steric screening, but previous work has shown with FTIR spectroscopy that in at least one case (PDIC+ED), the prevalence of these unreacted sites is minimal13. The surface site density is affected by steric screening or by the occurrence of inter-chain termination reactions, both of which reduce the number of reactive surface sites. When the total number of reactive sites decreases, less material can be deposited and the growth rate of the process decreases. The occurrence of termination reactions has been widely speculated to occur14–16,27,50–52, but little experimental

For the second category of effects that can influence MLD growth rate, i.e. film structure, the following factors play a role: 3) differences in film density and 4) changes in chain orientation. Differences in film density could explain changes in growth rate, since more tightly packed chains will take up less volume per mass deposited. However, given the observed 4x change in growth rates between the most flexible and most rigid chemistries, a similar 4x difference in density would be required to explain the difference in growth rate; such a density variation is not observed in X-ray reflectivity measurements, which show only a roughly 35% difference in density between the most and least flexible films (Supplemental Table 2). These densities are within the range of densities previously reported for similar systems53. As confirmation of this point, a recent study28 of the effect of aliphatic backbone flexibility on the growth rate and density of alucone MLD films observed a two-fold increase in growth rates between the shortest and longest precursor, with a only a 25% decrease in film density. Changes in chain orientation could also lead to different growth rates, since chains that grow more vertically will yield thicker films than those growing more horizontally, like with SAMs. Previous work has demonstrated that the growth rate of MLD films could be explained by the growth angle of the chemistry, which was determined using density functional theory calculations14,28; however, no previous work has examined this relationship experimentally. The observations of the present study allow us to discriminate between some of the above factors and their effect on the grow rates of the MLD films. Because we find no significant differences in the average bond orientation of the chains between the five backbone chemistries examined, despite the difference in growth rates, this suggests that the chain orientation relative to the surface (explanation 4) is not the main contributor to the observed growth rates of the chemistries studied here. Moreover, given that the chemistries exhibited saturation and that the density difference between the chemistries is minor, we can also rule out explanations 1 and 3. Hence, we conclude that the difference in growth rates between the MLD chemistries studied here is primarily the result of a reduced surface site density in the more flexible chemistries, explanation 2. Given the previously observed absence of unreacted sites in the PDIC+ED chemistry13, we propose that steric effects are a minor contributor to the reduction in growth rate


Page 11 of 17

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


and that the reduction in surface site density occurs by an increased frequency of double reaction terminations. This assertion is consistent with the observed data, in which films comprised of the more flexible monomers exhibit significantly lower steady state growth rates than those with more rigid monomers. However, a more explicit measurement of surface site density, such as an end group analysis, would provide further evidence of terminations. Double reaction terminations, in addition to explaining steady state growth rate differences between MLD chemistries, may also provide a rationale for the observed surface enhanced growth. In the first few cycles of growth, the number of surface sites has yet to be significantly reduced by terminations13. As the film becomes thicker, the total number of termination events has increased, reducing the overall growth rate of the chemistry. Support of this explanation can be seen through examination of the initial growth rates of the MLD processes, which are ~ 6-8 Å/cycle. Given that extended chain lengths of the chemistries are estimated to be between 12 Å and 17 Å (see above), an average chain orientation of 25° from the surface plane yields an estimated GPC of 5-7 Å/cycle. Hence the initial growth rates may be indicative of the inherent molecular length and orientation, before chain termination effects begin to play a major role. As the film gets thicker and more terminations occur, however, the growth rate begins to decrease until reaching a new steady state value. The reason that these terminations do not lead to a plateau in the film thickness as surface sites are consumed, and instead only reduce the growth rate to a new steady state value as observed here, is a discussion worthy of its own independent investigation. However, in brief: there are only a limited number of possible mechanisms that would allow the growth to continue: 1) the frequency of terminations becomes negligible once the number of surface sites is below a certain value, due to the low probability of termination occurring. This would lead to a slow, but perhaps negligible decrease in the growth rate with increasing cycles. Or 2) in some chemistries, like polyurea, precursors are able to absorb into the underlying film, either entangling with the underlying polymer or sticking to the surface through van der Waals forces or hydrogen bonding. This would increase the number of surface sites and could reach a steady state with the number of terminations taking place. Some sources have touched on this possibility51,52, but more work will be needed to deduce the mechanism behind this behavior.

In addition to answering questions about the MLD growth process, the current study allows us to probe the structure within the MLD films. One question about their structure is whether the chains (1) continue to grow at a particular angle set by their initial bonding to the surface, or if they (2) adopt multiple configurations with some overall average growth angle. The distinction between these is subtle, as both film structures could exhibit the same linear growth rate and the average bond angles in the chains. However, a film that includes chains with many twists, as might be found in a film with multiple configurations, is more likely to be amorphous, experience a higher degree of chain entanglement, and have a greater potential for the tuning of film properties by changing deposition conditions; hence, an understanding of the microscopic structure is important. A multiconfigurational growth model is also consistent with the idea that precursors can physisorb into the underlying film as described above, since these absorbed molecules would be unlikely to exhibit the same bond orientations as the rest of the film. On the other hand, a recent publication27 demonstrated growth of PDIC+PD films with what was believed to be a well-defined chain orientation. In that work, the authors used density functional theory to estimate a chain orientation and structure. They then compared this theory to experimentally measured values using a combination of ellipsometry and FTIR spectroscopy to determine chain orientation relative to the substrate and the orientation of the urea bonds, respectively. However, the authors’ observation of ordered growth does not exclude the possibility for multiconfigurational growth with similar systems or even with slight changes in the growth conditions. Although we cannot conclusively distinguish between these two models for the growth of polyurea MLD films, we suggest that multiconfigurational growth is more likely based on the presence of the out-of-plane ordering seen in the MLD systems studied here, as discussed below. Previous work27 suggested that the diffraction from the films was the result of stacked phenyl groups. However, in light of the fact that we observe this diffraction also with non-phenyl-based chemistries, we can determine that it doesn’t uniquely arise from phenyl stacking. Moreover, based on the strength of hydrogen bonding between urea groups54 and its common involvement in the underlying crystal structure of polyureas54, it seems more likely that the diffraction seen in the polyurea films grown here originates from the strong hydrogen bonding between urea groups in the chains. This assignment



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

Page 12 of 17

is further supported by the observed correlation between diffraction FWHM and the degree of order in hydrogen bonding (FTIR), i.e. the homogeneousbackboned films, which exhibited more ordered hydrogen bonding, show a greater degree of ordering. If the out-of-plane ordering is the result of hydrogen bonding, then it follows that at least some fraction of the urea groups are stacked on top of one another, forming layers of horizontally aligned polymer segments. Figure 7 illustrates some possible film structures. If the chains are tilted at 25° but exhibit purely outof-plane symmetry (Figure 7a), then the urea groups will not be aligned into their lowest energy state. Given the role hydrogen bonding plays in the structure of urea-based materials55, this structure is unlikely. If the chains are tilted and the urea groups are aligned (Figure 7b), the ordering of the chains will no longer be out-of-plane, and hence this structure is inconsistent with the experimental data. Only when the chains are horizontally aligned (Figure 7c) are they able to exhibit both out-ofplane ordering and alignment of the urea groups due to hydrogen bonding, making this structure the most likely origin of the out-of-plane ordering. However, because these chain lie at a 0o angle from the surface plane, in order to achieve an average chain tilt angle of around 25° (as measured in NEXAFS), some chains must also be more vertically aligned. This suggests a multiconfigurational growth model, with some fraction of chains adopting horizontal, paracrystalline orientations and some chains growing more vertically. The horizontal paracrystalline segments are observed primarily by diffraction, which would not be able to measure more amorphous vertically-aligned regions, while NEXAFS captures the average of the overall upward growth of the chemistry and the horizontal segments. The multiconfigurational growth is also captured by the more easily crystallizing BDIC+BD chemistry, which exhibits isotropic ordering in addition to its primarily outof-plane ordering. This implies that the chemistry is forming stacks of polymer chains in multiple directions.

Figure 7: Depictions of three possible chain orientations. The dashed lines represent the direction of repeat ordering in the structures. In (a), the urea groups and chains are tilted at 25° and have out-ofplane ordering, but poor hydrogen bonding due to a mismatch in the urea groups . In (b), the chains are also tilted at 25° and have good hydrogen bonding, but exhibit ordering that is not oriented directly out-ofplane. Only in (c) do the chains exhibit both good hydrogen bonding structure and out-of-plane ordering, which suggests that some chains must be horizontally aligned as they are in (c).


Page 13 of 17

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


The driving force behind the proposed formation of horizontally aligned chain segments is uncertain. However, hydrogen bonding between newly deposited monomers in each MLD cycle and the underlying surface can be reasonably suspected. For instance, studies of the adsorption of ethylenediamine on hydrogen bonding surfaces like silica suggest that monomers are energetically favored to orient face-on towards the underlying surface56. As coverage increases and available space near the surface decreases, new monomers may begin to attach more vertically, creating a mixture of horizontally and vertically aligned monomers in each layer with an average orientation of 25°. Chains with homogeneous backbone structures will likely have an easier time aligning into crystal structures with the underlying film, and the added flexibility of the aliphatic monomers makes it easier to bend into a horizontal arrangement. This picture agrees with the texturing of the out-of-plane peaks shown in Figure 4. The width of the diffraction peak becomes narrower with increasing backbone flexibility, which suggests that the most flexible backbones have an easier time aligning horizontally than do the more rigid backbones. A schematic representing the proposed growth behavior of these MLD films is summarized in Figure 8. According to the experimental data, in the first few cycles of growth, chains grow at an average angle of around 25° from the surface plane, leading to a film growth rate around 6-8 Å/cycle. As the film gets thicker, VASE data suggest that more

terminations take place, causing a slowing of the growth rate. MLD chemistries using the most flexible precursors exhibit more chain terminations, and consequently a lower growth rate. The average growth angle remains the same per the NEXAFS results, but the chains may adopt a variety of conformations as the film growth thicker, forming out-of-plane stacks of horizontally aligned segments consistent with GIXD results as well as vertically aligned domains. The spacing of these stacks and their relative crystallinities are dictated by the size of the monomers and the homogeneity of the resulting polymers. It is interesting to consider the application of these observations to hybrid organic-inorganic systems, since the presence of organic co-reactants is known22,37 to affect the growth rate and crystallinity of the resulting hybrid films compared to their purely inorganic counterparts. Given that these co-reactants are often used as a single layer between inorganic layers, these monomers are likely to adopt multiple orientations, resulting in a relatively unordered organic layer. This may affect the out-of-plane packing of the resulting films, as well as their electrical and thermal conductances. Longer, more flexible monomers have also been shown in one case28 to decrease the growth rate in a hybrid alucone system, which was similarly attributed to an increase in double-reaction terminations. Thus, study of the structure of the organic component in hybrid systems is worthy of further study.



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

Page 14 of 17

Figure 8: Depiction of thin film and bulk structure proposed for the films based upon the experimental results. In the first few cycles of growth, the average chain angle is 25° (NEXAFS) and few terminations have taken place (VASE). As the film gets thicker, terminations begin to reduce the growth rate (VASE) but the average growth angle remains the same (NEXAFS). Strong hydrogen bonding also causes some chains also align horizontally to form out-of-plane stacks of horizontally aligned polymer segments (FTIR, GIXD), forming a mixed paracrystalline and amorphous structure.

Conclusions In this work, we have presented results on MLD growth and film structure, based on experimental examination of the growth rate, crystallinity, and chain orientation of films grown with different backbone chemistries. Surprisingly, varying the backbone from a more rigid phenyl chemistry to a more flexible aliphatic chemistry does not have a strong impact on the resulting orientation of the chains, which all grow at an angle of ~ 25° from the surface plane, but results in a marked decrease in growth rate from 4 to 1 Å/cycle. Backbone choice also strongly affects crystallinity, since monomers with homogeneous backbones can be used to create films with much greater degrees of crystallinity. Despite previous assumptions that MLD films form relatively ordered layered structures, results from VASE, FTIR, GIXRD, and NEXAFS suggest that polymer chains in these films can adopt a number of different conformations, with some chain segments aligning horizontally with the surface to create diffuse out-of-plane paracrystallites. There is also strong evidence that MLD chemistries with more flexible backbones experience a greater number of

terminations than more rigid backbones, which reduces their overall growth rate.

Acknowledgements This work was supported by the National Science Foundation (CHE-1213879 and CHE-1607339). DSB would also like to acknowledge support from an NSF Graduate Research Fellowship. Part of this work was performed at the Stanford Nano Shared Facilities (SNSF) and the Stanford Synchrotron Radiation Lightsource (SSRL). Use of the SSRL at the SLAC National Accelerator Laboratory is supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC0276SF00515.

SUPPORTING INFORMATION AVAILABLE: MLD deposition conditions, NEXAFS measurement and correction details, AFM and XRR measurements used to verify film thicknesses, complete NEXAFS spectra for all chemistries at the carbon and nitrogen K-edges, graphitic carbon calibration standard, example NEXAFS fit, and peak fit parameters for all


Page 15 of 17

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


NEXAFS spectra. This material is available free of charge via the internet at http://pubs.acs.org.

(11)

Yoshimura, T.; Tatsuura, S.; Sotoyama, W. Polymer films formed with monolayer growth steps by molecular layer deposition. Appl. Phys. Lett. 1991, 59, 482 DOI: 10.1063/1.105415.

References (1)

Bergsman, D.; Zhou, H.; Bent, S. F. Molecular Layer Deposition of Nanoscale Organic Films for Nanoelectronics Applications. ECS Trans. 2014, 64, 87–96 DOI: 10.1149/06409.0087ecst.

(12)

Putkonen, M.; Harjuoja, J.; Sajavaara, T.; Niinistö, L. Atomic layer deposition of polyimide thin films. J. Mater. Chem. 2007, 17, 664–669 DOI: 10.1039/B612823H.

(2)

Zhou, H.; Bent, S. F. Molecular layer deposition of functional thin films for advanced lithographic patterning. ACS Appl. Mater. Interfaces 2011, 3, 505– 511 DOI: 10.1021/am1010805.

(13)

Kim, A.; Filler, M. A.; Kim, S.; Bent, S. F. Layer-bylayer growth on Ge(100) via spontaneous urea coupling reactions. J. Am. Chem. Soc. 2005, 127, 6123– 6132 DOI: 10.1021/ja042751x.

(3)

Zhou, H.; Blackwell, J. M.; Lee, H.-B.-R.; Bent, S. F. Highly sensitive, patternable organic films at the nanoscale made by bottom-up assembly. ACS Appl. Mater. Interfaces 2013, 5, 3691–3696 DOI: 10.1021/am4002887.

(14)

Loscutoff, P. P. W.; Zhou, H.; Clendenning, S. B. S. S. B.; Bent, S. F. S. Formation of organic nanoscale laminates and blends by molecular layer deposition. ACS Nano 2010, 4, 331–341 DOI: 10.1021/nn901013r.

(4)

Prasittichai, C.; Pickrahn, K. L.; Hashemi, F. S. M.; Bergsman, D. S.; Bent, S. F. Improving area-selective molecular layer deposition by selective SAM removal. ACS Appl. Mater. Interfaces 2014, 6, 17831– 17836 DOI: 10.1021/am504441e.

(15)

Adamczyk, N. M.; Dameron, a a; George, S. M. Molecular Layer Deposition of Poly( p -phenylene terephthalamide) Films Using Terephthaloyl Chloride and p -Phenylenediamine. Langmuir 2008, 24, 2081–2089 DOI: 10.1021/la7025279.

(5)

Ma, Y.; Martinez de la Hoz, J. M.; Angarita, I.; BerrioSanchez, J. M.; Benitez, L.; Seminario, J. M.; Son, S.B.; Lee, S.-H.; George, S. M.; Ban, C.; et al. Structure and Reactivity of Alucone-Coated Films on Si and LixSiy Surfaces. ACS Appl. Mater. Interfaces 2015, 7, 11948–11955 DOI: 10.1021/acsami.5b01917.

(16)

Du, Y.; George, S. M. Molecular Layer Deposition of Nylon 66 Films Examined Using in Situ FTIR Spectroscopy. J. Phys. Chem. C 2007, 111, 8509–8517 DOI: 10.1021/jp067041n.

(17)

Lee, J. S.; Lee, Y.-J.; Tae, E. L.; Park, Y. S.; Yoon, K. B. Synthesis of zeolite as ordered multicrystal arrays. Science 2003, 301, 818–821 DOI: 10.1126/science.1086441.

(18)

Loscutoff, P. W.; Lee, H.-B.-R.; Bent, S. F. Deposition of Ultrathin Polythiourea Films by Molecular Layer Deposition. Chem. Mater. 2010, 22, 5563–5569 DOI: 10.1021/cm1016239.

(6)

Piper, D. M.; Travis, J. J.; Young, M.; Son, S.; Kim, S. C.; Oh, K. H.; George, S. M.; Ban, C.; Lee, S.-H. Reversible high-capacity Si nanocomposite anodes for lithium-ion batteries enabled by molecular layer deposition. Adv. Mater. 2014, 26, 1596–1601 DOI: 10.1002/adma.201304714.

(7)

Luo, L.; Yang, H.; Yan, P.; Travis, J. J.; Lee, Y.; Liu, N.; Molina Piper, D.; Lee, S.-H.; Zhao, P.; George, S. M.; et al. Surface-Coating Regulated Lithiation Kinetics and Degradation in Silicon Nanowires for Lithium Ion Battery. ACS Nano 2015, 9, 5559–5566 DOI: 10.1021/acsnano.5b01681.

(19)

Yoshimura, T.; Ito, S.; Nakayama, T.; Matsumoto, K. Orientation-controlled molecule-by-molecule polymer wire growth by the carrier-gas-type organic chemical vapor deposition and the molecular layer deposition. Appl. Phys. Lett. 2007, 91, 33103 DOI: 10.1063/1.2754646.

(8)

Gould, T. D.; Izar, A.; Weimer, A. W.; Falconer, J. L.; Medlin, J. W. Stabilizing Ni Catalysts by Molecular Layer Deposition for Harsh, Dry Reforming Conditions. ACS Catal. 2014, 4, 2714–2717 DOI: 10.1021/cs500809w.

(20)

Ivanova, T. V.; Maydannik, P. S.; Cameron, D. C. Molecular layer deposition of polyethylene terephthalate thin films. J. Vac. Sci. Technol. A Vacuum, Surfaces, Film. 2012, 30, 01A121 DOI: 10.1116/1.3662846.

(9)

Liang, X.; Weimer, A. W. An overview of highly porous oxide films with tunable thickness prepared by molecular layer deposition. Curr. Opin. Solid State Mater. Sci. 2014, 19, 115–125 DOI: 10.1016/j.cossms.2014.08.002.

(21)

(10)

Zhou, H.; Bent, S. F. Fabrication of organic interfacial layers by molecular layer deposition: Present status and future opportunities. J. Vac. Sci. Technol. A Vacuum, Surfaces, Film. 2013, 31, 40801 DOI: 10.1116/1.4804609.

Atanasov, S. E.; Losego, M. D.; Gong, B.; Sachet, E.; Maria, J.-P.; Williams, P. S.; Parsons, G. N. Highly Conductive and Conformal Poly(3,4ethylenedioxythiophene) (PEDOT) Thin Films via Oxidative Molecular Layer Deposition. Chem. Mater. 2014, 26, 3471–3478 DOI: 10.1021/cm500825b.

(22)

Sundberg, P.; Karppinen, M. Organic and inorganicorganic thin film structures by molecular layer deposition: A review. Beilstein J. Nanotechnol. 2014, 5, 1104–1136 DOI: 10.3762/bjnano.5.123.



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

(23)

Léonard, S.; Harrowell, P. Macroscopic facilitation of glassy relaxation kinetics: Ultrastable glass films with frontlike thermal response. J. Chem. Phys. 2010, 133 DOI: 10.1063/1.3511721.

(24)

Dalal, S. S.; Walters, D. M.; Lyubimov, I.; de Pablo, J. J.; Ediger, M. D. Tunable molecular orientation and elevated thermal stability of vapor-deposited organic semiconductors. Proc. Natl. Acad. Sci. 2015, 112, 4227–4232 DOI: 10.1073/pnas.1421042112.

(25)

Liu, J.; Yoon, B.; Kuhlmann, E.; Tian, M.; Zhu, J.; George, S. M.; Lee, Y.-C.; Yang, R. Ultralow thermal conductivity of atomic/molecular layer-deposited hybrid organic-inorganic zincone thin films. Nano Lett. 2013, 13, 5594–5599 DOI: 10.1021/nl403244s.

(26)

DuMont, J. W.; George, S. M. Pyrolysis of Alucone Molecular Layer Deposition Films Studied Using In Situ Transmission Fourier Transform Infrared Spectroscopy. J. Phys. Chem. C 2015, 119, 14603–14612 DOI: 10.1021/jp512074n.

(27)

Park, Y.-S.; Choi, S.-E.; Kim, H.; Lee, J. S. FineTunable Absorption of Uniformly Aligned Polyurea Thin Films for Optical Filters Using Sequentially Self-Limited Molecular Layer Deposition. ACS Appl. Mater. Interfaces 2016, 8, 11788–11795 DOI: 10.1021/acsami.6b02142.

(28)

(29)

(30)

Park, Y.-S.; Kim, H.; Cho, B.; Lee, C.; Choi, S.-E.; Sung, M. M.; Lee, J. S. Intramolecular and Intermolecular Interactions in Hybrid Organic– Inorganic Alucone Films Grown by Molecular Layer Deposition. ACS Appl. Mater. Interfaces 2016, 8, 17489–17498 DOI: 10.1021/acsami.6b01856. Klepper, K. B.; Nilsen, O.; Francis, S.; Fjellvåg, H. Guidance of growth mode and structural character in organic-inorganic hybrid materials - a comparative study. Dalton Trans. 2014, 43, 3492– 3500 DOI: 10.1039/c3dt52391h. Nilsen, O.; Klepper, K.; Nielsen, H.; Fjellvåg, H. Deposition of Organic- Inorganic Hybrid Materials by Atomic Layer Deposition. ECS Trans. 2008, 16, 3– 14 DOI: 10.1149/1.2979975.

(31)

Klepper, K. B.; Nilsen, O.; Hansen, P.-A.; Fjellvåg, H. Atomic layer deposition of organic-inorganic hybrid materials based on saturated linear carboxylic acids. Dalton Trans. 2011, 40, 4636–4646 DOI: 10.1039/c0dt01716g.

(32)

Kubono, A.; Okui, N.; Tanaka, K.; Umemoto, S.; Sakai, T. Highly oriented polyamide thin films prepared by vapor deposition polymerization. Thin Solid Films 1991, 199, 385–393 DOI: 10.1016/00406090(91)90021-O.

(33)

Baker, J. L.; Jimison, L. H.; Mannsfeld, S.; Volkman, S.; Yin, S.; Subramanian, V.; Salleo, A.; Alivisatos, a P.; Toney, M. F. Quantification of thin film crystallographic orientation using X-ray diffraction with an area detector. Langmuir 2010, 26, 9146–9151 DOI: 10.1021/la904840q.

Page 16 of 17

(34)

Stöhr, J.; Outka, D. Determination of molecular orientations on surfaces from the angular dependence of near-edge x-ray-absorption finestructure spectra. Phys. Rev. B 1987, 36, 7891–7905 DOI: 10.1103/PhysRevB.36.7891.

(35)

Urquhart, S. G.; Hitchcock, A. P.; Priester, R. D.; Rightor, E. G. Analysis of polyurethanes using core excitation spectroscopy. Part II: Inner shell spectra of ether, urea and carbamate model compounds. J. Polym. Sci. Part B Polym. Phys. 1995, 33, 1603–1620 DOI: 10.1002/polb.1995.090331105.

(36)

Rubinstein, M.; Colby, R. H. Polymer Physics; Oxford University Press: New York, 2006.

(37)

Lee, B. H.; Yoon, B.; Abdulagatov, A. I.; Hall, R. A.; George, S. M. Growth and properties of hybrid organic-inorganic metalcone films using molecular layer deposition techniques. Adv. Funct. Mater. 2013, 23, 532–546 DOI: 10.1002/adfm.201200370.

(38)

Ramin, M. a.; Le Bourdon, G.; Daugey, N.; Bennetau, B.; Vellutini, L.; Buffeteau, T. PM-IRRAS investigation of self-assembled monolayers grafted onto SiO 2/Au substrates. Langmuir 2011, 27, 6076– 6084 DOI: 10.1021/la2006293.

(39)

Ramin, M. a.; Le Bourdon, G.; Heuzé, K.; Degueil, M.; Buffeteau, T.; Bennetau, B.; Vellutini, L. EpoxyTerminated Self-Assembled Monolayers Containing Internal Urea or Amide Groups. Langmuir 2015, 31, 2783–2789 DOI: 10.1021/la5049375.

(40)

Coleman, M. M.; Sobkowiak, M.; Pehlert, G. J.; Painter, P. C.; Iqbal, T. Infrared temperature studies of a simple polyurea. Macromol. Chem. Phys. 1997, 198, 117–136 DOI: 10.1002/macp.1997.021980110.

(41)

Jimison, L. H.; Himmelberger, S.; Duong, D. T.; Rivnay, J.; Toney, M. F.; Salleo, A. Vertical confinement and interface effects on the microstructure and charge transport of P3HT thin films. J. Polym. Sci. Part B Polym. Phys. 2013, 51, 611– 620 DOI: 10.1002/polb.23265.

(42)

Kasai, N.; Kakudo, M.; Kasai, N.; Eroglu S., Toprak S., Urgan O, MD, Ozge E. Onur, MD, Arzu Denizbasi, MD, Haldun Akoglu, MD, Cigdem Ozpolat, MD, Ebru Akoglu, M. X-ray diffraction by macromolecules; Springer Berlin Heidelberg: New York, 2005; Vol. 33.

(43)

Sirringhaus, H.; Brown, P. J.; Friend, R. H.; Nielsen, M. M.; Bechgaard, K.; Langeveld-Voss, B. M. W.; Spiering, a. J. H.; Janssen, R. a. J.; Meijer, E. W.; Herwig, P.; et al. Two-dimensional charge transport in self-organized, high-mobility conjugated polymers. Nature 1999, 401, 685–688 DOI: 10.1038/44359.

(44)

Hattori, T.; Takahashi, Y.; Iijima, M.; Fukada, E. Piezoelectric and ferroelectric properties of aliphatic polyurea films synthesized by vapor deposition polymerization. In 9th International Symposium on Electrets (ISE 9) Proceedings; IEEE, 1996; Vol. 2199,


Page 17 of 17

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


Oligomer orientation in vapor-molecular-layerdeposited alkyl-aromatic polyamide films. Langmuir 2012, 28, 10464–10470 DOI: 10.1021/la3017936.

pp 819–824. (45)

(46)

(47)

(48)

Ishihara, H.; Kimura, I.; Yoshihara, N. Studies on segmented polyurethane—urea elastomers: Structure of segmented polyurethane—urea based on poly(tetramethylene glycol), 4,4’diphenylmethane diisocyanate, and 4,4’diaminodiphenylmethane. J. Macromol. Sci. Part B 2006, 22, 713–733 DOI: 10.1080/00222348308245751. Born, L.; Hespe, H. On the physical crosslinking of amine-extended polyurethane urea elastomers: A crystallographic analysis of bis-urea from diphenyl methane-4-isocyanate and 1,4-butane diamine. Colloid Polym. Sci. 1985, 263, 335–341 DOI: 10.1007/BF01412250. Hattori, T.; Iijima, M.; Takahashi, Y.; Fukada, E.; Suzuki, Y.; Kakimoto, M. A.; Imai, Y. Synthesis of aliphatic polyurea films by vapor deposition polymerization and their piezoelectric properties. Japanese J. Appl. Physics, Part 1 Regul. Pap. Short Notes Rev. Pap. 1994, 33, 4647–4651 DOI: 10.1143/JJAP.33.4647. Hattori, T.; Iijima, M.; Takahashi, Y.; Fukada, E.; Suzuki, Y.; Kakimoto, M. A.; Imai, Y. Piezoelectric properties of thin films of aliphatic polyurea 79 synthesized by vapor deposition polymerization. Ferroelectrics 1995, 171, 363–372 DOI: 10.1080/00150199508018447.

(49)

Eda, K.; Okazaki, T.; Yamamura, K.; Hashimoto, M. Three-dimensional supramolecular assembly having infinite two-dimensional interlocking networks built up only from simple and non-rigid organic molecules via hydrogen bonds. Crystal structures of α,ω-diureidoalkanes H2N(CO)NH–(CH2)n– NH(CO)NH2 with n=4 and. J. Mol. Struct. 2005, 752, 93–97 DOI: 10.1016/j.molstruc.2005.05.042.

(50)

Peng, Q.; Efimenko, K.; Genzer, J.; Parsons, G. N.

(51)

Dameron, A. A.; Seghete, D.; Burton, B. B.; Davidson, S. D.; Cavanagh, A. S.; Bertrand, J. A.; George, S. M. Molecular Layer Deposition of Alucone Polymer Films Using Trimethylaluminum and Ethylene Glycol. Chem. Mater. 2008, 20, 3315–3326 DOI: 10.1021/cm7032977.

(52)

Seghete, D.; Hall, R. A.; Yoon, B.; George, S. M. Importance of trimethylaluminum diffusion in three-step ABC molecular layer deposition using trimethylaluminum, ethanolamine, and maleic anhydride. Langmuir 2010, 26, 19045–19051 DOI: 10.1021/la102649x.

(53)

Zhou, H.; Toney, M. F.; Bent, S. F. Cross-Linked Ultrathin Polyurea Films via Molecular Layer Deposition. Macromolecules 2013, 46, 5638–5643 DOI: 10.1021/ma400998m.

(54)

Steed, J. W.; Atwood, J. L. Supramolecular Chemistry; John Wiley & Sons, Ltd: Chichester, UK, 2009.

(55)

Rotello, V. M.; Thayumanavan, S. Molecular Recognition and Polymers: Control of Polymer Structure and Self-Assembly; Rotello, V. M., Thayumanavan, S., Eds.; John Wiley & Sons, Inc.: Hoboken, NJ, USA, 2008.

(56)

Xu, M.; Liu, D.; Allen, H. C. Ethylenediamine at air/liquid and air/silica interfaces: Protonation versus hydrogen bonding investigated by sum frequency generation spectroscopy. Environ. Sci. Technol. 2006, 40, 1566–1572 DOI: 10.1021/es051537l.

Table of Contents Figure


Effect of Backbone Chemistry on the Structure of Polyurea Films

Recommend Documents