To appear in: “Chemical Applications of Synchrotron Radiation”, T. K. Sham, Ed. World Scientific Publishing Co. Ltd. Singapore

NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

H. Adea and S.G. Urquharta,b

a Department of Physics, North Carolina State University, Raleigh, NC 27695

harald_ade@ncsu.edu

b (As of Jan. 2000) Department of Chemistry, University of Saskatchewan, 110 Science Place, Saskatoon, SK S7N 5C9 Canada

stephen.urquhart@usask.ca

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Table of contents 1. Introduction 2. NEXAFS Spectroscopy

2.1. Instrument Aspects of NEXAFS and Electron Energy Loss Spectroscopy 2.2. NEXAFS Spectroscopy of Polymers: Chemical Sensitivity

2.2.1. Chemical Sensitivity: Examples 2.3. NEXAFS Spectroscopy of Polymers: Orientation Sensitivity 2.4. NEXAFS Spectroscopy of Polymers: Quantitative Analysis

2.4.1. Blends of Polycarbonate and Poly(butyl terephthalate) 2.4.2. Quantitative Analysis of the Chemical Composition of Polyurethanes

3. NEXAFS microscopy: Instrumentation and Analysis Tools 3.1. Quantitative Image Analysis

4. NEXAFS Microscopy: Applications 4.1. Surfaces and Thin Films

4.1.1 Polyimide Films and Surfaces 4.1.2. Overcoats and Lubricating Layers on Hard Disks 4.1.3. Confined, Free Standing Homopolymer Thin Films 4.1.4. Thin Film Blends 4.1.5. Thin Film Polymer Bilayers 4.1.6 Block Copolymer Thin Films

4.2. Bulk Morphology and Composition of Multicomponent, Multiphasic Polymers 4.2.1. NEXAFS Microscopy of Poly(ethylene terephthalate)/ VectraTM Blends 4.2.2. Engineered Polyurethane Polymers 4.2.3. Multilayers 4.2.4. Elastomer Composites 4.3. Fibers 4.4. Organic Geochemistry 4.5. Meteoritics and Interplanetary Dust Particles 4.6 Soils and Environmental Studies 4.7. Other Applications

5. Summary and Future Outlook Acknowledgements Tables Figure Captions References

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1. Introduction Carbonaceous materials, such as synthetic and natural (including

biological) polymers, exhibit a rich carbon, nitrogen and oxygen K-edge Near Edge X-ray Absorption Fine Structure (NEXAFS). The spectral variations that are observed might even constitute the most complex set of NEXAFS features to be found and acquired from any class of materials. This is primarily due to the essentially covalent nature of the chemical bonds in these materials and the numerous possible combinations of single and multiple bonds between carbon atoms as well as hetero-atoms (oxygen, nitrogen, chlorine, fluorine, sulphur, etc.). Simple building blocks can react to result in complex macromolecules with millions of atoms. Furthermore, many polymeric systems, such as engineered polymers, latexes, multilayers, organic geochemical, and meteoritic materials, etc., are naturally, accidentally or intentionally in-homogeneous or structured. The combination of the chemical complexity of constituent polymers and the microheterogeneity found in many materials results in a multitude of systems that can benefit from micro/nano NEXAFS analysis. It is, thus, not surprising that the characterization of the surface and the bulk of polymeric materials has been one of the most frequent application of NEXAFS microscopy techniques. We will first exemplify the basic underlying power of NEXAFS microscopy by discussing the NEXAFS spectra from a wide variety of polymers and small molecule model analogs. We will subsequently review the majority of NEXAFS microscopy applications of polymeric systems to date, spanning disciplines ranging from Polymer Science, Organic Geochemistry, Meteoritics to Environmental Studies. We will only briefly mention biological x-ray microscopy applications, even though all biological systems are made from “natural polymers”. Our intent is to provide a broad-based tutorial for readers unfamiliar with the field, rather than to provide an exhaustive summary for experts. 2. NEXAFS spectroscopy

The chemical sensitivity of NEXAFS spectroscopy is at the heart of the unique capabilities of x-ray microscopy for polymer microanalysis. In x-ray microscopy, NEXAFS spectroscopic transitions can be used as a chemically sensitive image contrast mechanism or micro-NEXAFS spectra can be extracted from small sample volumes. The range and subtlety of x-ray microscopy applications ultimately depends on our comprehension of the “chemical information content” or the “structure-spectral relationships” of NEXAFS spectroscopy.

NEXAFS spectroscopy measures the photoabsorption cross-section for the excitation or photoionization of tightly bound core electrons. These spectra are element specific, as each element has a characteristic core binding energy (i.e. Carbon 1s: ∼290 eV, Nitrogen 1s: ∼400 eV, Oxygen 1s: ∼530 eV, etc). The spectral features correspond to transitions from the ground state to a core excited state. In general this must be treated as a multielectron process. However, it is convenient to describe x-ray absorption in closed-shell molecules with an orbital approximation as a one-electron transition (C 1s → π*)

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perturbed by the creation of the core hole. Figure 1 presents a schematic of the x-ray photoabsorption process for the Carbon 1s NEXAFS spectrum of poly(styrene-r-acrylonitrile). This NEXAFS spectrum is dominated by intense and narrow features at low energy (285-288 eV). The first transition at 285 eV is a C 1s → π* electronic transition of the phenyl functional group and the second transition at 287 eV is a C 1s → π* electronic transition of the acrylonitrile functional group. The energy of these features is dictated by the combination of initial state effects (core binding energies) and final state effects (energy of the optical orbital). At higher energy, broad C 1s → σ* transitions are superimposed on the photoionization continuum. NEXAFS spectroscopy has been reviewed by Stöhr1, and complementary, but somewhat limited, overviews of polymer NEXAFS spectra have been presented by Kikuma2, Unger3 and Ade4,5.

We will first discuss instrumentation aspects of polymer NEXAFS spectroscopy and the analogous Electron Energy Loss Spectroscopy (EELS) technique (§2.1). The nature of the chemical sensitivity (§2.2) and the orientational sensitivity (§2.3) of polymer NEXAFS spectroscopy will follow. The use of NEXAFS in x-ray microscopy for quantitative chemical analysis is subsequently presented and discussed (§2.4). 2.1. Instrument Aspects of NEXAFS and Electron Energy Loss Spectroscopy

A NEXAFS spectrum can be recorded by scanning the incident x-ray photon energy across a core absorption edge and measuring the probability for photon absorption by the sample. This process can be measured indirectly through the emission of secondary electrons (partial or total yield) or fluorescent photons (fluorescent yield)1 or directly by measurement of the photon transmission through the sample.

Complementary core excitation spectra can also be acquired by Electron Energy Loss spectroscopy (EELS)6,7. In EELS, inelastic scattering of a monoenergetic electron beam induces electronic excitation and ionization in the material. The “energy loss” spectrum of this electron beam contains the core excitation spectrum. Most of the early polymer core excitation spectroscopy literature, particularly at the C 1s core edge, was performed by EELS spectroscopy8-11. EELS spectroscopy can be performed in a TEM microscope12-14 and is analogous to performing NEXAFS spectroscopy on small sample volumes. Energy filtered TEM imaging15 is an analogue to imaging by NEXAFS microscopy. The relatively poor spectral resolution in commercial instruments (discussed below) and higher beam damage limits the use of the EELS methods for compositional analysis of complex polymeric materials.

Figure 2 highlights how spectral resolution affects the chemical sensitivity of the spectroscopy by comparing the C 1s spectrum of poly(ethylene terephthalate) (PET) recorded at different energy resolutions by

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x-ray absorption and by electron energy loss techniques16-18. The fact that higher energy resolution increases the chemical sensitivity of a spectroscopy is hardly surprising, but this figure illustrates how much spectral resolution is useful. The lower resolution C 1s EELS spectra of PET (recorded using a LaB6 thermionic emission electron source or a cold field emission gun) show that PET is unsaturated from the presence of the C 1s → π*C=C transition at 285 eV. Higher resolution spectra can resolve the carbonyl C 1s → π*C=O transition at 288 eV. The bottom trace, recorded with ~0.15 eV Full Width Half Maximum (FWHM) energy resolution not only resolves the aromatic C 1s → π*C=C transition and the splitting in this transition, but also small bands above the C 1s → π*C=O transition at 287 eV are now clearly resolved18. Recently, vibronic features were identified in the NEXAFS spectra of polystyrene isotopomers recorded with 50 meV resolution19. An energy resolution of better than 100 meV is required to realize all the subtleties in the NEXAFS spectra of polymers, while an energy resolution of less than 50 meV will most likely not reveal additional features. While x-ray absorption spectroscopy was used to acquire the higher resolution spectrum of PET presented in Figure 2, we note that EELS spectroscopy is capable of similar or higher energy resolution spectroscopy. Fink et al. has acquired C 1s EELS spectra of several polymers with an energy resolution better than 0.2 eV FWHM8,9. While field emission microscopes are capable of ~0.5 eV FWHM energy resolution, “descan” procedures used to minimize radiation damage can degrade the instrumental resolution to more than 1 eV. Specialized TEM-EELS instruments have demonstrated 0.2 - 0.4 eV FWHM energy resolution for semiconductors20 and nickel aluminum intermetallics21 and the development of TEM-monochromators capable of reaching 50 meV resolution is now in progress22,23. 2.2. NEXAFS Spectroscopy of Polymers: Chemical Sensitivity

This section will discuss the chemical and functional group sensitivity of the NEXAFS spectroscopy of polymers. We will first outline the general characteristics of the NEXAFS spectra of unsaturated and saturated polymers, followed by a discussion of the models and approaches used to understand and interpret these spectra.

The vocabulary of molecular orbital (MO) theory will be used to describe the polymer NEXAFS spectra since it is appropriate for discussing the spectral-structural relationships present in covalently bonded materials. Figure 3 presents the C 1s NEXAFS spectra for a range of polymers containing unsaturated functional groups (phenyl rings and double bonds). The NEXAFS spectra of these unsaturated polymers are dominated by low energy C 1s → π*C=C transitions at ~285 eV. The shape of this π*C=C band varies with the chemical and electronic structure of these unsaturated polymers. An example of this chemical sensitivity can be observed in the C 1s NEXAFS spectra of the polyurea and polyurethane polymers (left panel of figure 3). In these spectra,

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the core → π*(LUMO) transition has two major components, C 1s(C-H) → π *C=C component at ~285 eV, and a C 1s(C-R) → π*C=C component at ~286.5 eV. Note that in this notation, the core level is indicated parenthetically before the arrow (i.e. C-H), and the nature of the upper level of a given transition is indicated by the final subscript (i.e. C=C for phenyl). In the spectra of TDI polyurea and TDI polyurethane the C 1s(C-H) → π *C=C transition is split into two narrowly spaced, relatively unresolved spectral features. The peak at ~286.5 eV corresponds to core excitations from the “C-R” phenyl carbon atoms that are attached to the amide groups. The inductive effect of the amine group shifts the C 1s ionization potential of the C-R phenyl carbon to lower energy, increasing the energy of the C 1s(C-R) → π *C=C transition to the common manifold of π*(LUMO) states. In the MDI polyurea and polyurethane spectra, the C-R peak is about half as intense as in the TDI polyurea and polyurethane spectra. This reflects the higher fraction of C-R(amide) bonds per phenyl ring in the TDI polymers relative to the MDI polymers.

Figure 4 presents the C 1s NEXAFS spectra of a selection of saturated polymers (polyolefin and polyethers). The NEXAFS features in these spectra are predominantly core → σ* electronic transitions with perhaps some Rydberg character24, although these are expected to be attenuated in the solid state25. In Figure 4, the C 1s NEXAFS spectra of polyethylene (PE), polypropylene (PP), polyisobutylene (PIB) and ethylene propylene rubber (EPR) differ slightly. The main chemical differences between these polymers are the addition of methyl groups to the (CH2)x backbone. In the spectrum of PIB, a low energy shoulder is more intense as there are two –CH3 groups in the PIB repeat unit, relative to one in PP and none in PE. The C 1s spectra of polyethylene oxide (PEO) and polypropylene oxide (PPO) are shifted to higher energy by the core binding energy shift induced by the oxygen atom. The methyl (-CH3) group contribution in PPO is observed as a low energy shoulder at ~287 eV, at approximately the same energy as similar features of other polymer carbon atoms that are not directly bonded to oxygen (PP, PIB).

Several related concepts are used somewhat interchangeably to discuss NEXAFS spectra: “building blocks”, “functional group fingerprinting” and “molecular modeling” approaches. The “building block model” is a semi-quantitative approach, where the spectra of building blocks (diatomic or functional group fragment, etc.) are used to simulate the NEXAFS spectra of more complicated species. “Functional group fingerprinting” is a qualitative approach, where the unique spectroscopic signatures of chemical units present in polymers (e.g. phenyl, carbonyl, amide, etc) are established using suitable spectroscopic models. In “molecular modeling”, the spectra of small molecules (for which a considerable literature exists26) is used to establish these qualitative fingerprints. Specific considerations for the use of NEXAFS spectroscopy for quantitative analysis will be discussed below (§2.4).

The chemical and functional group sensitivity of NEXAFS spectroscopy has been discussed in terms of the spectroscopic signature of “diatomic” building blocks1. This description originates in valence bond theory, where

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chemical properties are described in terms of the additivity of diatomic bond properties. In the approach described by Stöhr1, the spectra of polyatomic molecules can be viewed as an assembly of diatomic or pseudodiatomic “building block” contributions. For example, the C 1s spectrum of acetonitrile (CH3CN) has been viewed as the sum of CN, C-C and C-H “diatomic” contributions (i.e. core → σ*CN, σ*C-C and σ*C-H transitions)1. However, bond/bond interactions (hyperconjugation and electronic delocalization) will reduce the “diatomic” character of σ* core excited state in molecules that are more complex than diatomics. If the σ* MO is not localized to specific diatomic bonding pairs, then the core → σ* transitions will not reflect diatomic bonding character in an additive fashion. These conjugation and delocalization effects are one origin of breakdowns of the bond length correlation model, which attempted to relate the term value (energy relative to the ionization potential) of σ* core excited states with the bond length of a particular bond27,28.

For organic molecules and polymers it is useful to consider larger building blocks such as functional groups. For example, polymers containing phenyl groups share the same C 1s(C-H) → π*C=C transition as benzene29,30. The use of spectral signatures of larger functional group “building blocks” (i.e. phenyl, ketones, esters, etc.) to interpret polymer spectra makes both chemical and electronic sense. The bond/bond interactions that are absent in the “diatomic fragments” will be present in suitable polyatomic models. For example, in a phenyl building block, the effects of cyclic π-delocalization are incorporated where they would be absent in a pseudodiatomic building block models (i.e. ethylene). Petterson et al. enumerated several difficulties and limitations of the building block approach. In particular, the geometry of the experimental or calculated building block model can deviate from the geometry of the polymer, and the building blocks lack interconnecting bonds present in the polymer31.

The concept of a “critical size” for building block models is illustrated in Figure 5. If PET is viewed as a sum of phenyl and ester functional groups (concentrating here on the unsaturated species), then its spectrum may be modeled by the sum of the spectrum of benzene and formic acid32. As seen in Figure 5, this model reflects the gross character of the spectrum of PET (low energy aromatic C 1s → π*C=C and the characteristic carbonyl C 1s → π*C=O transition), but the energy of the carbonyl C 1s → π*C=O is slightly shifted in energy. The molecule ethyl benzoate could be chosen to more accurately model the electronic delocalization between the phenyl and ester group in PET. This model is slightly more effective, but fine details present in the spectrum of PET are still not well reproduced. The model that best reproduces the spectrum of PET is dimethyl terephthalate, which reproduces the split C 1s → π*C=C transition (two peaks resolved in PET), as well as the small features above the carbonyl C 1s → π*C=O transition. The spatial extent of electronic delocalization (the substitution of two ester groups on the phenyl ring) is essential for

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reproducing the features of this spectrum. This demonstrates that an interpretive model for the NEXAFS spectrum of PET must include the relevant spatial extent of electronic delocalization to provide an accurate functional group fingerprint.

In polymers with extensive electronic delocalization, such simple and relatively localized functional group or building block descriptions may not be appropriate. In a theoretical study of polyene oligomers with different chain lengths, Carravetta et al. predicted a strong decrease in the π* intensity with increasing chain length, suggesting that π-electron screening decreased the effect of core hole localization33. This effect can not be explained by the building block model. 2.2.1. Chemical Sensitivity: Examples

Figure 6 compares the NEXAFS spectra of MDI-polyurethane models34 (denoted BM, B2 and B0, with structures indicated in Fig. 6) with the previously reported gas-phase Inner Shell Electron Energy Loss Spectra (ISEELS) spectrum of ethyl N-phenyl urethane35 and the solid phase NEXAFS spectrum of 2-propane diol. These MDI polymers have different proportions of what have been called “hard segment” (aromatic rich) and “soft segment” (polyether rich) components (i.e. two different building blocks). The B0 polymer has a high polyether concentration while the B2 polymer has a smaller polyether concentration and the BM polymer is entirely hard-segment. Ethyl N-phenyl urethane and 2-propane diol are used as molecular models for the urethane and polyol chemical structures present in the MDI polyurethanes (BM, B2 and B0). The model ethyl N-phenyl urethane is excellent at reproducing the spectroscopic features of BM, including the intense C 1s(C-H) → π*C=C transition at ~285 eV, the weaker C 1s(C-R) → π*C=C transition at ~286.5 eV, and the sharp C 1s(C=O) → π*C=O transition at 290 eV. Ethyl N-phenyl urethane structurally reproduces the electronic delocalization between the phenyl and carbonyl groups, and it differs from the polymer at the aliphatic linkages (which should have a minimal effect on the unsaturated contribution to the spectrum). The main difference between the polymer BM and the molecular model is the spectral resolution: ~0.7 eV FWHM for ISEELS; 0.2 eV FWHM for NEXAFS.

While the aromatic component of these polymers is well modeled by ethyl N-phenyl urethane, the limitations of 2-propane diol as a fingerprint model for poly(propylene oxide) have been explored by Urquhart et al34. Figure 7 presents the C 1s NEXAFS of poly(propylene oxide) (PPO) and solid-phase 2-propane diol, and the gas phase C 1s ISEELS spectrum of 2-propane diol. The C 1s spectra of 2-propane diol is used as a molecular model for the C 1s spectrum of PPO as this molecule reproduces the nearest neighbor-bonding environment present in PPO. In aliphatic molecules and polymers such as PPO, the local bonding environment primarily determines the NEXAFS spectrum. Here, the enhanced Rydberg contribution in the gas phase spectrum of 2-propanol creates a significant difference relative to the spectrum of PPO. In

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addition, the energy of the C 1s → σ* transition (at ~293 eV) is shifted between the polymer and the molecular models, suggesting the effect of σ-delocalization in the polymer spectra34. Therefore, both phase and bonding effects must be considered when using molecular model spectra to interpret polymer spectra.

A complementary study to the BO, BM, B2 polyurethanes mentioned above explored the applicability of the building block approach for biological molecules with the aim to predict bio-polymers such as proteins and peptides from amino acid building blocks36. Carbon NEXAFS spectra from six amino acid monomers, four dipeptides and one tripeptide were measured and compared. The formation of peptide bonds was judged to have a relatively small influence on the NEXAFS spectra, so that the weighted sum of monomer spectra closely approximates the spectrum of the small peptides (see Fig. 8). These results open up the possibility of predicting NEXAFS spectra of proteins from their amino acid content. Although averaging over many amino acids may reduce the effectiveness of NEXAFS contrast, in some cases it may be possible to use the inherent differences in the NEXAFS spectra of different proteins to map and identify regions of high concentration of a particular protein37. Theoretical calculations and a comparison to the experimental amino acid spectra have been performed by Carravetta et al. 38. Pioneering NEXAFS spectroscopy studies of biopolymers have been carried out by Kirtly et al. who have characterized the electronic environment of nitrogen in nucleic acid bases, nucleotides, polynucleotides and DNA39.

Additional examples where the chemical sensitivity of NEXAFS spectroscopy was used to characterize the electronic and/or geometric structure of organic and related materials include: i.) Photoactive polymers: poly(phenylenevinylene)40, poly-p-phenylenes and

polyacenes41, photo-oxidation of electroluminescent polymers42, and UV polymerized diacetylene43

ii.) Plasma damaged polymers: oxygen plasma damaged polypropylene44 and argon plasma damaged polycarbonate45

iii.) Metal coated polymers: chromium coated poly(ethylene terephthalate)46 and metalized polycarbonate47

iv.) Electron and x-ray beam damaged or decomposing polymers: hexadecanethiolate monolayers48 and PMMA49-51

v.) Organic geochemical materials: Petroleum Asphaltenes and Coals52,53 and Kerogens and Bitumens54

vi.) Inorganic polymers: polydimethylsilane films55,56 and polydi-N-hexyldisiloxane57.

2.3. NEXAFS Spectroscopy of Polymers: Orientational Sensitivity

Spectral features in the NEXAFS spectroscopy of aligned molecules and polymers exhibit linear dichroism. Specifically, their intensity depends on the

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angle between the x-ray electric field vector eE and the transition moment vector, <Ψf|e••r|Ψi>. The electric field vector eE of synchrotron radiation produced by bending magnets and most undulator devices is linearly polarized in the plane of the storage ring. The transition intensity varies as:

θcos22

E | r e | I ∝•∝ where θ is the angle between the electric-dipole vector r and the electric-field vector eE: The intensity of the transition is largest when the electric field vector (eE) of the radiation is aligned with the transition moment vector (eE • r = 1).

This effect is best demonstrated by example. Figure 9 presents the angle-dependant C 1s NEXAFS spectrum of highly oriented polytetrafluoroethylene (PTFE) thin films recorded by Ziegler et al.58 When the electric field vector eE is directed perpendicular to the polymer chain (α = 90o), the intensity of transitions to σ*C-F states (labeled “1” in Figure 8) are enhanced relative to σ*C-

C states (2), since the orientation of the σ*C-F electric-dipole vector is oriented normal to the polymer chain. When the electric field vector eE is directed along the polymer chain (α = 0o), the relative intensity of these transitions is reversed. Determining the absolute polymer orientation from the angle-dependant NEXAFS spectra is simpler for planar & conjugated polymers than for aliphatic polymers. In planar conjugated polymers there is clear σ/π separation, so the absolute orientation of the π* transitions is unambiguously known. In aliphatic polymers (e.g. polytetrafluoroethylene and polyethylene), the absolute orientation of the electronic transition moment is a matter of interpretation. Is the orientation of the C 1s → σ*C-C resonance aligned along the polymer axis or along the individual C-C bonds? Originally, Stöhr and Outka proposed the alignment of the σ*C-C resonance along the C-C bond to interpret their results59 and recent Xα calculations of propane were used by Väterlein et al. to reaffirm

this conclusion60. In contrast, ab inito calculations of heptane were employed by Hähner et al.61 to argue that the σ* C-C resonance is aligned along the polymer axis.

Notwithstanding this issue, orientational studies are one of the largest areas of application of NEXAFS spectroscopy to polymers. These studies include various forms of polyethylene (CH2)x, such as oriented homopolymer59,62; oligomers such as hexatriacontane62 and pentacontane63 and Langmuir-Blodgett films 61,64. Other examples include polytetrafluororethylene homopolymer and models58,62,65-67; polyethylene terephthalate and related polymers (polybutylene terephthalate and polyethylene naphthalate)68-70; self-assembled monolayers, including n-alkanethiols71, semifluorinated amidethiols72; and electrochemically polymerized thin films such as poly(3-methyl thiophene)73. This is not a comprehensive list, but a indication of the wide range of applications.

A recent area of interest has been studies of surface versus bulk

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orientation in polymers. Here, the depth sensitivity of different detection methods (Auger, total electron yield and/or fluorescence) is used in combination with linear dichroism spectroscopy to examine differences in orientation. This effect has been used to study surface alignment in spun-cast polystyrene films74,75 and the surface orientation in rubbed polyimides76-79. We will present examples of the use of NEXAFS orientational dependence for polymer characterization by x-ray microscopy below. 2.4. NEXAFS Spectroscopy of Polymers: Quantitative Analysis

In transmission NEXAFS Spectroscopy, optical density spectra (OD) can be derived via Lambert-Beer’s law from the transmitted X-ray intensity as:

where an energy scan from the sample (I) is normalized to another energy scan recorded without a sample (I0). Quantitative spectral analysis is then provided by the dependence of the absorbance:

where µµ is the energy dependent mass absorption coefficient, ρρ is the density, and t is the sample thickness.

For accurate quantitation, well-characterized NEXAFS spectra of carefully chosen analytical models of the polymer components are required. A basic assumption for quantitative analysis is that the spectra of non-interacting moieties or components are additive. If a mixed system contains x% of A and y% of B, and A and B do not interact electronically, then the composition of the system can be determined by fitting linear combinations of the pure model spectra of A and B to the mixed AB system. For a blend of two or more homopolymers (i.e. polystyrene/poly(methyl methacrylate)), the analytical models can simply be the individual homopolymers. For quantitation of components in a random block copolymer (i.e. poly(styrene-r-acrylonitrile)), the spectra of the homopolymers (polystyrene and polyacrylonitrile) can be used, provided that the blocks do not interact electronically. We will now outline two examples that detail the power and limitations of using NEXAFS spectroscopy for quantitative analysis: a binary mixture of two partially miscible polymers (polycarbonate and poly(butylene terephthalate)) and the quantitative analysis of a complex polyurethane. The utility of quantitative analysis by micro-NEXAFS spectroscopy in the x-ray microscope will be demonstrated below. 2.4.1 Blends of Polycarbonate and Poly(butyl terephthalate) Polycarbonate (PC) and poly(butylene terephthalate) (PBT) are partially miscible polymers. Depending on composition and temperature, blends of PBT and PC exhibit a two-phase morphology with a mixed PC/PBT composition. NEXAFS microscopy is useful for determining the composition of these phases. Figure 10 presents the average C 1s NEXAFS spectrum acquired from a PC/PBT blend, in comparison to the C 1s spectra of the PC and PBT

( )IoIOD ln−=

( )IoItODA ln−=== µρ

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homopolymers. By least-squares analysis of the mixed polymer spectrum or principal component analysis80 the composition of the polymer blend can be determined. The example presented yielded an accuracy of 1% of the expected composition from blend formulation. 2.4.2. Quantitative Analysis of the Chemical Composition of Polyurethanes

Polyurethanes are complex materials that may consist of a variety of component species, depending on both the reagents and the conditions of polymerization81. The goal in the study of these materials is to quantify and map the composition, in particular the urea and urethane content, and possibly that of minority species, in a variety of formulations and thereby determine relationships between chemical composition, processing, and properties of the final polymer product. Spatially resolved quantitative analysis of functional group composition (particularly the urea and urethane content) is needed to understand the chemical basis of microstructure formation in polyurethane polymers81-83. In order to determine the quantitative analysis abilities of x-ray microscopy for these materials, a spectroscopic study was performed on test polyurethane polymers, where a reliable estimate of ether, urea and urethane content could be made a priori from the reaction stoichiometry.

Figure 11 (left) presents the C 1s spectra of the carbamate model (C: TDI-polyurethane), the urea model (U: TDI-polyurea) and the polyether model (E polyether-polyol) used as standards for the quantitative analysis using fitting procedures described below. A detailed analysis of these spectra is presented elsewhere84. The right of Figure 11 shows a close-up view of the C 1s spectra of the three test polymers (codes 258, 259, 260), along with the best quality fits. The spectra of the test polymers (258, 259, 260) and the stripped spectra of the individual analytical components (C, U, E) were background subtracted and normalized in the far continuum on an oscillator strength per atom basis. The C, U and E composition of 258, 259 and 260 was determined by a linear least squares fit over the energy range 282-286 eV and 289-291 eV. This choice of energy range restricts the fit to the energy range that is most sensitive to the chemical differences and where the models are most representative of the chemistry. The results of the NEXAFS analysis are compared to the composition predicted from the formulation chemistry in Table 1.

The quality of the match between data and optimized model (Fig. 11) demonstrates that C 1s NEXAFS spectra, when analyzed using spectra of appropriate models recorded with the same experimental conditions, can determine chemical composition at the ~10 mol% level, with generally a few % accuracy (only two data points in Table 1 are off by more than 10%). Good energy resolution (<0.2 eV, preferably <0.1 eV) is a critical factor in being able to track the subtle changes in the line shapes in the 289-291 eV range which provide the sensitivity to quantitative composition. Energy scale accuracy and stability is also important, as model and analyte spectra must be placed on a common energy scale. Another important factor is confidence in the absolute spectral shape for both the models and the analyte materials. Changes in energy

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resolution, linear dichroism in oriented samples, higher-order x-ray photons as well as detector dark noise and nonlinearity are all factors that potentially can affect the fidelity of spectroscopic features and thus the accuracy of NEXAFS quantitation.

3. NEXAFS microscopy: Instrumentation and Analysis Tools

Given the wealth of spectroscopic detail exhibited by NEXAFS spectra, it is not surprising that NEXAFS microscopy is developing into a useful research tool for the micro-characterization of polymers and other organic materials. Still, sensitivity to specific moieties and functional groups can in many or even most cases be exceeded by infra red, NMR and Raman spectroscopies81. The power of NEXAFS microscopy is thus mostly based on its ability to exceed the spatial resolution of these other spectroscopy techniques. To date, NEXAFS microscopy with a spatial resolution of better than 50 nm has been accomplished both in transmission to measure bulk properties, and in a reflection geometry to study surfaces. This level of spatial resolution is at least an order of magnitude better than what can be accomplished with complementary techniques. Furthermore, exceptional surface sensitivity of a few nanometers can be achieved in NEXAFS microscopy. We will briefly discuss some of the most salient features of the various technological approaches to NEXAFS microscopy in the 200 to 900 eV range that are presently being employed.

High spatial resolution instruments for transmission experiments typically employ zone plate optics. The accuracy of the zone placement and the width of the outermost zone determines the spatial resolution achievable with these elements. Presently, zone plates with outermost zone widths between 35 and 80 nm are in use. Two types of instruments have been developed and are being continually refined: i) Scanning Transmission X-ray Microscopes (STXM) that typically generate a microprobe about 50 nm in diameter and mechanically raster scan the sample85,86, and ii) full-field conventional transmission x-ray microscopes (TXM), in which a zone plate downstream of the sample magnifies the sample onto a 2-D detector87-89 (See schematics in Fig. 12). Besides high spatial resolution, NEXAFS microscopy with good chemical sensitivity also requires good energy resolution. For STXM and the 200 to 900 eV energy range this is achieved with a grazing incidence monochromator upstream of the microscope. A resolving power of several thousand is routinely obtained. Near the carbon edge, the most widely used edge for carbonaceous materials, a resolving power of 2000-5000 is readily available86,90,91. TXMs presently employ zone plate monochromators in order to match the numerical aperture of the illumination at the sample to the numerical aperture of the objective. Unfortunately, present zone plate monochromators used in a TXM have relatively low energy resolution (∼0.5 eV), limiting the utility of TXM for NEXAFS microscopy of polymers.

A technical challenge with zone plate optics results from the fact that the focal length is proportional to the photon energy. This requires continual

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refocusing by changing the sample to zone plate distance during a photon energy scan of a point spectrum. Care and effort must be exercised to control the transverse motion of the scanned element (sample or zone plate) during refocusing to avoid blurring of the region from which the spectrum is to be acquired. This places severe constraints on the accuracy of the mechanical apparatus, an accuracy that has not yet been consistently achieved in existing STXMs. One way to work around the accuracy problem, at least partially, is to acquire image sequences at each photon energy in the spectrum92. The images in these sequences can than be correlated to each other and shifted to provide the proper alignment. If one is only interested in spectra from a few locations, the image sequence approach results in a time penalty: Sequences typically require between 0.5-3 hours, while a few point spectra require a few minutes. However, the relative certainty and control over the sampling area are typically well worth the extra time. It is also more straightforward with image sequences to keep the sample in focus. Refocusing after a change in energy is also necessary in a TXM, with similar considerations regarding stability and alignment. In a TXM, one has to translate both the zone plate objective and the detector in order to keep the magnification the same.

STXMs are presently only operated at undulator beamlines91,93,94, although the first STXM had been operated at a bending magnet beamline of a NSLS UV ring. The requirement to illuminate the zone plate coherently and to provide a high data rate is most easily accomplished at the brightest sources. In order to increase STXM capacity at relatively low cost, an effort is underway to build a STXM at a relatively bright bending magnet at the ALS. The instrument and beamline will be optimized for C-NEXAFS. The brightness requirements for TXMs are less severe and TXMs are very successfully operated at bending magnet sources even at second generation synchrotron facilities. While there are several efforts to build x-ray microscopes based on various laboratory laser-based sources, we believe that the requirements imposed by the spectroscopy will make laboratory based NEXAFS microscopy a relatively distant possibility in the future. To study surfaces, it is necessary to be able to detect secondary processes that occur after the absorption of a photon, processes that track the absorption cross section as a function of photon energy relatively well. Fluorescent photons and secondary or primary photoelectrons fulfill that need. All these secondary processes are extensively utilized in NEXAFS spectroscopy without spatial resolution, in what is called the fluorescent yield (FY), total- (TEY) or partial electron yield (PEY) mode, respectively1. In NEXAFS microscopy, “PEY” mode is most readily accomplished in a Photoemission Electron Microscope (PEEM) as the PEEM instrument itself acts as a low pass energy filter with a few eV bandwidth (traditionally, “PEY” refers to experiments utilizing a bias voltage on the detector to only accept electrons with high kinetic energy, i.e. in a high pass filter mode, in order to achieve high surface sensitivity). The photoelectrons emitted from the sample are accelerated by an electric field in a cathode lens and magnified onto a 2-dimentional detector with

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the use of electrostatic or electromagnetic lenses (see Fig. 13 as an example). Contrast in PEEM is not only determined by composition, i.e. the cross section of the material at a given photon energy, but also by topography and illumination. Topographic contrast may arise from the reduced or enhanced collection efficiency in the PEEM from locations that are not perpendicular to the optical axis. Viewed from the traveling direction of the electrons, concave round features tend to have reduced collection efficiency, while convex features are enhanced. Sharp features tend to have enhanced electric fields and therefore might have enhanced emission. In addition, the sample is illuminated at a shallow angle of 30° or less and the local photon intensity of the sample depends on the topography. Partial or complete shadowing as well as increased emission can contribute to contrast in PEEM images. While the overall collection efficiency of the photoelectrons and the illumination contrast is effected by topography, spectra can still be acquired from any location of the sample that is at least partially illuminated. Unless the topographic contrast or charging depends on the photon energy, and hence would result in artifacts, any spectrum can in principle be normalized to the pre- and post-absorption edge continuum and quantitative analysis on a per atom basis can be performed with PEEM spectra. Since PEEMs acquire information in parallel from all pixels in the image, and incoherent illumination of the sample is possible, the source requirements are generally not extreme. PEEMs can be successfully operated at any bending magnet beamline that provides sufficient energy resolution95-97. PEEMs are also relatively easy to operate and do not require any refocusing during photon energy changes. However, for future ultra high spatial resolution (<10nm) instruments, the necessary flux densities on the sample can only be provided with undulator beamlines. 3.1. Quantitative Image Analysis

We previously discussed how a building block approach can be utilized to quantitate the composition by simulating a given spectrum with a linear combination of spectra from model polymers. If whole images are to be analyzed quantitatively, a Singular Value Decomposition (SVD) procedure is often used98-101. SVD requires prior knowledge about the sample and works only if the composition of the sample is known. In addition, the linear absorption coefficients of the components must be accurately known in order to minimize systematic errors in the analysis. The NEXAFS spectra used as standards for SVD must be free of spectral distortions caused by saturation due to the presence of higher order spectral contamination, detector non-linearity102 or the use of different energy resolving power between the different experiments that might effect the peak heights.

The SVD procedure requires a series of images acquired at a number of energies that equal or exceed the number of compositional components. The reference spectra provide crucial guidance for the selection of photon energy. Usually energies are chosen that correspond to large spectral differences

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between the constituent components. We show an example of raw transmission images in Fig. 14. Each transmission image is converted to an optical density (OD) image by computing OD=-ln(I/I0), where I is the transmitted intensity and I0 is the incident intensity. The normalization with I0 can be accomplished most readily if the image itself contains an open area (see Fig. 15), but recording a separate I0 spectrum through an open area of the sample is also possible. Since OD=αt=µρt, where α is the linear absorption coefficient, t the sample thickness, µ is the mass absorption coefficient and ρ is the density, each pixel in an OD image can be described by a linear combination of the product of the linear absorption coefficient α and the thickness t for each component respectively. This results in a set of linear equations that can be solved for the thickness of each component if the same or more images than components have been acquired. Since SVD utilizes only a limited and small number of images, it results in the smallest (and minimum) dose to the sample to extract required information. Since αt=µρt, one could alternatively use the mass absorption coefficient as the known quantity and calculate the mass thickness ρt, rather than the thickness itself. The latter is advantageous if the same component might be present in the sample with different densities, i.e. a polymer in an amorphous and crystalline form. The relative contrast and intensities of the images resulting from the SVD analysis are not affected by this change in parameters, simply the interpretation of these images.

4. NEXAFS Microscopy: Applications

Numerous NEXAFS microscopy applications of carbonaceous materials have been performed during the last couple of years in a variety of fields ranging from polymer science to meteoritics. We will structure our discussion along a loose organizing principle that progresses from the “simple” to the apparently or supposedly “more complex”. We thus first describe applications to homopolymer surfaces, then single or multicomponent thin polymer films, bulk properties of multicomponent polymers, polymer fibers, organic geochemical materials (coal, coke), meteorites and organic soils in environmental studies. We complement the spectromicroscopy applications with a short discussion of other x-ray applications involving carbonaceous, including biological, materials. 4.1. Surfaces and Thin Films

Thin polymer films have considerable technological importance and are used in numerous applications ranging from multi-color photographic printing, paints, adhesives and flat panel and liquid crystal displays, to dielectric insulators in electronic devices, membranes and protective and lubricating coatings. Understanding thin film characteristics, including the influence of the interfaces on the thin film properties and the interface properties themselves, is crucial for a rational improvement of the performance of polymers in these applications. However, fundamental aspects of the behavior of polymers in reduced-dimensional systems and surfaces/interfaces are still too poorly

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understood to aid improvements effectively. Thin polymer films and surfaces thus enjoy continued research interest. NEXAFS microscopy and spectroscopy provide a promising approach to study a variety of scientific and technological issues, such as, for example, the phase separation process of thin film blends. NEXAFS microscopy has also been utilized to investigate the stability and dewetting behavior of bilayers and to asses the damage caused to lubricating films by wear in electronic storage systems (see sections below).

NEXAFS spectroscopy by itself can and has been utilized to study certain surface properties and characteristics of polymers due to its excellent surface sensitivity (about 2-10 nm, rather material dependent) and its polarization dependence. In fact, the potential for NEXAFS characterization of surfaces is enormous. A large number of polymer surface characterizations, e.g. in adhesion science, are currently performed with XPS, even though NEXAFS spectroscopy would be much more suitable to the task. Some recent, non-exhaustive examples of NEXAFS spectroscopy applications are the investigation of the alignment and orientation of polymers and its functional groups in polyimide surfaces after rubbing with a buffing cloth76,77, the surface relaxation of PS74, and the orientation of semi-fluorinated block copolymers103. Additional examples have been mentioned previously in section 2.2.

4.1.1 Polyimide Films and Surfaces

Thin films of polyimide are extensively utilized in flat panel displays. In these devices, polyimide films about 0.2 µm thick are typically coated onto the electrode surfaces. Buffing the surface with a velour cloth is the traditional means to prepare a surface that aids the alignment of the liquid crystal (LC) molecules in the final device. Understanding the conditions and factors that cause LC alignment and how that is related to possible polymer alignment and functional group orientation in the substrate provides valuable information about the function and the possible improvement of flat panel displays. In order to investigate relevant parameters, Cossey-Favre et al. utilized NEXAFS microscopy in a PEEM to characterize the inhomogeneity of the surface orientation produced with rubbing the polyimide with a velour cloth104. It was clearly observed that the rubbing process does not orient the surface of the film homogeneously. Stripes based on orientation contrast several microns in widths were observed in the polyimide film. However, polyimide films produced under the same rubbing conditions resulted in LC films with the liquid crystal oriented uniformly over the entire sample surface. Since the separation distance between the oriented stripes of polymers are less than the correlation length of the liquid crystal, the heterogeneous distribution of the pinning centers is sufficient to align the liquid crystal and the oriented polymer areas must act as pinning centers to align the liquid crystal.

NEXAFS microscopy also allowed the measurement of the orientation of polymer chains of BPDA-PDA polyimide surfaces in response to rubbing with a stylus of well defined shape and with a known load. The stylizing process is

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an ideal model case to correlate the polymer orientation with the known local stress, a parameter that can only be estimated poorly for buffing with a cloth. An ellipsoidal stress profile was produced on the polymer surface by utilizing a stylus with a radial tip. NEXAFS analysis with a PEEM was used to map the polymer response. We show an example of one of these tracks produced by rubbing a stylus across a polyimide in Fig. 16. Spectra inside the track from two different sample orientations relative to the polarization orientation of the x-rays are shown in Fig. 17. The minimum force necessary to cause alignment was determined by measuring the width of the track that resulted in aligned molecules and comparing it to the stress profile. The minimum normal stress necessary to orient the surface of the BPDA-PDA film was found to be 45 MPa, much lower than the bulk yield stress of 200-300 MPa. The lower value at the surface was attributed to a reduction of polymer entanglement at the surface.

4.1.2. Overcoats and Lubricating Layers on Hard Disks

NEXAFS microscopy with the PEEM also proved very successful in studying the tribo-chemical processes between the slider of a hard disk drive, the hard carbon overcoat and the lubricant95,105,106. Wear tracks on a supersmooth disk coated with 5 nm cathodic arc deposited hard carbon and lubricated with 0.85 nm of perfluoropolyether are shown in Fig. 18(a). The scratch in the middle of the micrograph is caused by one of the rails of the slider. The carbon and fluorine NEXAFS spectra from the undamaged area, the wear track caused by the rail, and the area between rails are shown in figures 18(b) and (c). The spectra in the undamaged area and between the rails are identical. In the scratch caused by one of the rails, the fluorine is almost completely removed and the carbon spectrum shows a new peak at 288.5 eV, attributed to the formation of carboxylic bonds. Anders et al. also compared the wear tracks created with two kinds of sliders: uncoated sliders (Al2O3/TiC) and sliders coated with sputter deposited CHX 95,106. A comparison of the tracks created by the coated and uncoated sliders showed that the modification of the lubricant and the formation of the carboxyl bond is stronger for the uncoated slider than for the coated slider. This was attributed to a catalytic reaction between the lubricant and the Al2O3/TiC of the uncoated slider. The corresponding surfaces on the sliders themselves were also examined. The NEXAFS spectra from inside and outside wear marks on the sliders have shown considerable differences. The uncoated slider had tracks containing significant amounts of carbon with a spectrum that was very similar to that observed at the edge of the wear track on the disk. Both the carbon and oxygen K-edge NEXAFS spectra showed evidence for carboxyl groups. The coated slider had undamaged CHx in most areas, but the scratches that were observed on the slider itself had the same characteristic peak at 288.5 eV observed from other damaged areas.

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4.1.3. Confined Free-standing Homopolymer Thin Films Patterning of thin polymer films is attracting considerable and increasing interest, both from a fundamental polymer science perspective as well as for technological applications107. Formation of holes can take place in free standing films, driven by dispersion forces which amplify structural instabilities at the film surface. Mechanical confinement of polymer films offers the possibility to control these instabilities, and raises the possibility to generate unique self-assembled structures in polymeric systems. Thin continuous layers of silicon oxide (SiOx) confining a very thin (nominally ~50 nm) polystyrene film prevent hole formation. However aggressive annealing at temperatures well above the PS glass transition temperature generates an in-plane micro-structure due to the attractive dispersion force between the SiOx-air surface108. These structures provide a model-system for understanding and developing control mechanisms for one of the fundamental forces governing self-assembly in polymer systems.

To completely model the mechanisms for this spontaneously generated lateral morphology, it is important to measure the thickness of the trapped polymer film throughout the structure. The optical density obtained from quantitative NEXAFS microscopy readily provides this information. It had been assumed that there was almost complete exclusion of the polystyrene from the thin, non-continuous regions. However, NEXAFS microscopy indicated ~30% residual PS in the thinnest parts of the sample (see Fig. 19)109, a fact that now needs to be incorporated into the models.

4.1.4. Thin Film Blends

Generally, less is known about the properties of polymer blends when they are processed into thin films than for bulk blends. For example, numerous studies110-119 have shown that for binary polymer blends the spinodal decomposition and coarsening process can be more complex in thin films due to the various boundary conditions that are imposed. Interfacial energies are important since they determine thermodynamic equilibrium structures. The viscosity near interfaces and surfaces can be modified, which will influence the dynamics, and possibly cause the formation of kinetically trapped morphologies. A number of morphologies that evolve during the phase separation process in polymer thin films have been previously observed112-118,120. However, the composition of two and three-dimensional morphologies and their time dependent evolution has never been directly determined, particularly in the early stages of phase separation. NEXAFS microscopy provides a new approach to study such systems100,121-123.

The first NEXAFS microscopy feasibility experiments on phase separated thin film blends were performed by Smith et al. 124,125 and Cossy-Favre et al. 121. Smith et al. observed good chemical contrast between PS and PBrS phases in phase separating blends of these polymers, as well as indications that the PS might encapsulate the PBrS during the late stages of

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annealing. Cossy-Favre et al. observed that differences in composition in phase separated PS and polyvinylmethylether (PVME) and PS/poly(styrene acrylonitrile) blends can be detected from the surfaces of thin films that had been annealed for 3 hrs at 170°. This study also showed that the formation of PS protrusions in the PS/PVME film is not accompanied by a total dewetting of the polymers from the substrate.

The first NEXAFS microscopy studies of thin polymer films explicitly exploring the time evolution and composition of domains focused on the morphologies formed during spincasting and annealing of polystyrene (PS) and brominated polystyrene (PBrS) blends. Emphasis was placed on the late stage morphology of thin films of PS and PBrS, after annealing above the glass transition temperature of both components122. Quantitative NEXAFS maps (SVD) were compared to AFM topographs and complemented with Secondary Ion Mass Spectroscopy (SIMS) depth profiling. The morphologies observed depended strongly on blend composition. NEXAFS microscopy of annealed samples showed directly that the morphology changes from droplets to surface holes in a continuous PBrS layer as the PBrS concentration increases and when PBrS becomes the majority phase. SIMS data indicated continuous PS at the substrate and the surface interface. The encapsulation provides for a continuous PS phase for all blend compositions, and explained the observed “droplet” and “hole” structures which are formed for different PBrS fractions. Although the interfacial energy could have been minimized if the PS formed spheres in the PBrS-rich phases, it is the continuity constraint of the PS at the interfaces that requires the formation of the hole morphologies. From the morphologies formed, constraints for the values of the various interfacial energies could be deduced. (Below, we will show that PEEM microscopy can provide a much more direct and unambiguous means of showing PS encapsulation of PBrS than the combination of area averaged SIMS, surface AFM, and bulk STXM measurements.)

PS/PMMA blends are another excellent model system to investigate the dynamics and the morphologies formed during phase separation in thin films. The combination of quantitative NEXAFS mapping via SVD methods and lateral force microscopy allowed the observation and characterization of a variety of new phenomena and morphologies during both the early and late stages of the phase separation process100. The model polymer blends utilized were monodisperse, 50/50 weight percent PS/PMMA blends spun cast out of toluene solution onto Si substrates. The 143 nm thick films were annealed at elevated temperatures in a vacuum, quenched to room temperature, and subsequently characterized with NEXAFS microscopy. An example of the morphologies observed as a function of annealing time is presented in Fig. 20. Quantitative NEXAFS microscopy determined the composition of the mixed phases that resulted from solvent spin casting. The sudden rearrangement into domains much smaller than those originally formed was observed after very short annealing times. Subsequently, unique, jagged hydrodynamic mass flow patterns formed during the early stages of coarsening which are in qualitative

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agreement with recent simulations of phase segregation in two-dimensional viscous fluids126. Complicated polymer-polymer interfaces persisted even in the later stages that were explained in terms of the geometric constraints of a thin film and the dependence of polymer viscosity on film thickness. The observations of the early stages are particularly interesting because the lattice Boltzmann simulations by Wagner and Yeomans126 on two-dimensional viscous binary fluids indicated that the early stages of the spinodal decomposition process are even more sensitive to the dimensionality of the film than the later stages. These results predict that the dominant growth mechanism in thin films depends on the competition between the relative viscosities, the diffusivities of the two liquids, as well as the amplitude of the surface capillary waves relative to the domain size. In contrast to bulk spinodal decomposition, the resulting patterns are predicted to lack scale-invariance and to develop with time-dependent rate laws. Depending on the viscosities, different morphologies, including relatively jagged ones, can evolve. The unstable domains observed with NEXAFS microscopy are even more jagged than those of the model. A 50 nm thick PS/PMMA (27k/27k film on Si annealed at 165°C, and a 50 nm thick 50/50 weight percent blend of 90k PS and 27k PMMA on Au showed the same kind of instability. In contrast, a 50 nm layer of the 90k/27k blends on Si did not show these instabilities123. As can be seen from the above example, polymer blends that are processed into thin films can produce surface modulations with amplitudes much larger than the original film thickness due to the phase separation between the immiscible polymers (see for example Fig. 20R). This presents a serious problem in many applications. To improve the miscibility of polymer blends, copolymer compatibilizers can be added, which will reduce the interfacial tension if the compatibilizer is located at the polymer/polymer interface. This process is typically limited by the tendency of the compatibilizer to form micelles in one of the phases, and for strongly segregated systems the formation of a microemulsion has never been achieved. In order to probe the influence of confinement on the process of compatibilization Zhu et al. investigated PS/PMMA thin films that incorporated a PS-b-PMMA diblock copolymer127. Samples of a bottom PMMA layer and a top layer of a blend of PS and 30% of a PS-b-PMMA diblock copolymer were annealed and investigated with NEXAFS microscopy and scanning force microscopy. When the system size is comparable to the size of the bulk micelles, complete miscibility is achieved, and the resulting two dimensional microemulsion is stable. Specifically, as the PS/PS-b-PMMA layer thickness was reduced to less than the size of the bulk micelle diameter (<50nm), the two homopolymers completely mixed and formed a microemulsion. Compositional NEXAFS mapping was used to determine that the phases in the microemulsion extended all the way through the thin film, and also to observe the micelles that form in the PS phase for larger film thicknesses. Thus, Zhu et al. observed that the loss in configurational entropy as a result of the confinement can change the micellar transition for copolymers. These findings are important as mixing in

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thin films can be achieved independently of the specific polymer chemistry by making the film thin enough. 4.1.5. Thin Film Polymer Bilayers Additional NEXAFS investigations of PBrS/PS model systems focused on the dynamics and morphology formation during the dewetting of a bilayer consisting of a 50 nm thick PBrS film on top of a 30 nm thick PS film99. These investigations were the first combination of surface (PEEM) and quantitative bulk (STXM) characterization of a polymeric system to assess the three dimensional morphology. As the PBrS is dewetting the PS sublayer, holes form randomly and subsequently grow to form Veronoi tesselation patterns. These patterns consist of an interconnected network of spines that eventually break up to form droplets. The thickness maps of the constituent polymers PBrS and PS, as well as the total thickness maps of the bilayer annealed for 1 week, are shown in Figure 21. It is easy to see from the line profiles that the spines consist of sharply delineated PBrS and are at least partially wetted, if not encapsulated, by PS walls. PEEM studies of the same type of sample provided surface NEXAFS spectra from the top 10 nm of the surface from a variety of sample areas including the spines and intersection of the spines. The PEEM results did not detect any spectral feature at 286 eV, an energy characteristic of the C-Br shifted C 1s to π* transition, from any of the sample areas (see Figure 22). Therefore, the PBrS spines are completely encapsulated. Characterization of the time evolution - from 0 to 11 days of annealing - of the dewetting process with NEXAFS spectroscopy and PEEM furthermore revealed the encapsulation pathway and allowed to distinguish whether diffusion or flow of the polymers are the dominant processes. The results indicated that the encapsulation is delayed until a sufficient number of holes have formed in the PBrS layer through which the PS can penetrate the PBrS film and subsequently flow along the PBrS/air and PBrS/PS interfaces. Additional studies showed that the apparent contact angle at the polymer air interface decreases exponentially with film thickness with a constant that is dependent upon the radius of gyration Rg (a measure of the size of a random coil polymer that increases with the molecular weight of the polymer)128. NEXAFS data show that the droplets consist of a PBrS core fully encapsulated by PS for substrate thicknesses greater than Rg while only partial encapsulation is seen for substrates with thickness less than Rg. 4.1.6 Block Copolymer Thin Films

Block copolymers have been extensively investigated using a wide variety of analytical techniques including electron microscopy, x-ray scattering, and optical microscopy. These techniques yield considerable information about the morphology and organization of these materials. NEXAFS microscopy is a key complementary tool providing unique information. For example, Carl Zimba et al. investigated the novel molecular organization of a high molecular weight rod-coil block copolymer129. This block copolymer has a polystyrene

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(PS) coil block of 9K molecular weight and a main-chain poly(hexyl isocyanate) (PHIC) rod block of 245K molecular weight. Homopolymer PHIC is known to behave as a lyotropic nematic liquid crystal in a wide range of solvents with a persistence length of 50-60 nm. Rotational motion along the backbone of the PHIC is limited by short-range interactions, leading to inherent chain stiffness. The PS behaves as a typical coil polymer having a low persistence length. Combined as a block copolymer, there is competition between the microphase separation of the PS and PHIC blocks into periodic structures and the tendency of the PHIC block to form anisotropically ordered structures, giving rise to new morphologies130. Thin film samples were cast from 0.05% solutions in toluene onto both uncoated and carbon-coated TEM grids. The carbon-coated grids were used for transmission electron microscopy (TEM) and selected area electron diffraction (SAED) while the uncoated grids were used for NEXAFS microscopy. A typical image obtained using TEM is shown in Fig. 23(a) and clearly shows a microphase separated morphology with long-range order over tens of microns. The PS domains, stained with RuO4, appear as dark regions shaped somewhat like arrowheads while the PHIC regions appear as largely featureless light regions between rows of PS domains. An interesting feature in this image is the orientation of the PS arrowheads which changes direction between adjacent PS-rich layers. Obtained from a sample area of 10-20 lamellae, the SAED pattern (inset of Fig. 23(a)) shows a superposition of two distinct single crystal-like PHIC patterns rotated relative to each other by approximately 90 degrees, indicating that the PHIC chain axis alternates between +45 degrees and -45 degrees with respect to the geometric normal of the lamellae layers. Using an x-ray energy of 285.0 eV, corresponding to the phenyl carbon resonance of the PS, the NEXAFS image (Figure 23(b)) clearly shows a similar arrowhead morphology as the TEM image, albeit at lower spatial resolution. At 288.5 eV, which corresponds to the carbonyl resonance of the PHIC, NEXAFS images reveal the structure of the PHIC layers in ways not available with TEM (Fig. 23(c)). In Fig. 23(c), the PHIC domains appear as layers of alternating intensity separated by the PS domains that appear as white pearls. The intensity variation of the PHIC layers arises from differences in dichroic absorbance of the polarized synchrotron radiation due to orientation of the carbonyl moiety, and thus also the backbone, of the PHIC. In contrast to the SAED pattern, the NEXAFS image is able to give detailed information about the orientation of individual PHIC domains. In fact, NEXAFS microscopy reveals previously unsuspected micro-structure within the PHIC domains. 4.2. Bulk Morphology and Composition of Multicomponent, Multiphasic Polymers Many polymeric materials of academic and industrial interest are multicomponent polymers that have multiple phases in the bulk. Control over the composition and size distribution of the domains and their interface properties is often very important in determining the materials properties. A variety of blends and multiphasic polymers have already been investigated with

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NEXAFS microscopy. These include studies of: the morphology of poly(ethylene terephthalate) (PET) -polycarbonate (PC)131, LDPE-PET-Kraton132, and polycarbonate-ABS blends132; the morphology and composition of PET – oxybenzoate/oxynapthoate (VectraTM ) blends133, rubber

toughened poly(methyl methacrylate) (PMMA)134, and macrophase-separated random block copolymer/homopolymer blends 135 and the characterization of phase separation during processing, such as precipitates in polyurethanes 5,84,136 and multi-phase liquid crystalline polyesters132 and single- and multi-step synthesized latexes109. We include here as examples a discussion of the characterization of PET/VectraTM blends, various polyurethane polymers, multilayers and rubber composites. 4.2.1. NEXAFS Microscopy of Poly(ethylene terephthalate)/ VectraTM Blends

The morphology and composition of poly(ethylene terephthalate) (PET) and VectraTM A950, (73/27 mol% oxybenzoate/ 2,6oxynaphthoate) blends produced by mechanical alloying has been investigated with NEXAFS microscopy by Smith et al. 133. PET is a widely used, semi-crystalline, thermoplastic polymer, while VectraTM is a nematic liquid crystalline polymer (LCP) with excellent barrier properties. Since incorporation of LCPs into commodity polymers remains an ongoing challenge in the design of new high-performance, low-cost polymeric blends, the novel method of mechanical alloying has been used to create blends of a thermoplastic polymer and a nematic LCP. No preferential, heavy metal stain exists for enhancing the contrast between these two polymers in Electron Microscopy and spectroscopic means had to be utilized to differentiate these two components in the blends. Although relatively similar functional groups, such as aromatic and carbonyl groups, are present in each polymer, the NEXAFS spectra of each polymer are quite distinct. Reference spectra and the chemical structure of these polymers are shown in Fig. 2. NEXAFS microscopy had little difficulty delineating the morphology in these materials. Melt-pressed films of PET/VectraTM blends varying in composition — (A) 75/25, (B) 90/10 and (C) 99/1 — are displayed for comparison in Fig. 24. The 75/25 blend has been milled for 24 hrs at –180°C, whereas the other two blends were cryogenically milled for 6 hrs. After milling, the powders of each blend were melt-pressed at 285°C (above the melting points of PET and VectraTM) for 5 min and cryo-microtomed at –100°C to obtain sufficiently thin sections for STXM. All of the micrographs in Fig. 24 were acquired at 286.7 eV and show Vectra (dark) domains dispersed throughout a PET matrix. These dispersions exhibited a broad size distribution, ranging from 100 nm to about 20 µm in diameter. The frequency and size of the Vectra dispersions increased with increasing Vectra concentration. The presence of these smallest Vectra domains demonstrated that mechanical alloying is capable of pulverizing Vectra. In addition, these studies showed that the Pet/VectraTM blends retain much of the degree of mixing imparted by alloying after post-processing in the molten state, and that the VectraTM

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dispersions contain little, if any PET. Molecular orientation of the VectraTM domains was investigated with linear dichroism microscopy. Figs. 25A and B show the same sample region imaged with a different orientation relative to the linearly polarized x-ray beam by rotating the sample by 90°. The images have been converted to optical density for analysis, hence the inverted contrast relative to Fig. 24. Fig. 25C shows the dichroic ratio image, mapping the degree of orientation in this material. Medium gray corresponds to a dichroic ratio of zero and hence no molecular orientation. Anisotropic orientation of the VectraTM molecules was only observed above the signal to noise in domains larger than 2 µm. 4.2.2. Engineered Polyurethane Polymers Engineered polymers constitute a versatile and ubiquitous class of materials of enormous economic value and continued development potential. Polyurethanes alone are a US$2 billion dollar business with major applications in automotive and home furnishings. In the past several years, demand for improved polyurethane performance has skyrocketed as businesses strive to introduce new products and compete in a global environment. Despite its ubiquitous utilization and considerable past research, several fundamental aspects of polyurethanes are not well understood. In general, the distribution of the so-called hard segment urea is believed to have a major influence on the mechanical properties of the foam (modulus, compression set, load bearing, etc.), but rather little is known beyond the average size and distances of the hard-segment urea as measured with x-ray and neutron scattering techniques. Another key parameter influencing physical properties in the foams is the morphology and distribution of additive or modifier phases. These phases are smaller than a few microns in size since larger particles can be stress concentrators which diminish performance. The size scale inherent to these issues creates challenges for traditional characterization techniques since the size range is beyond that of common spectroscopic methods for chemical speciation (IR, NMR). Through several studies over the past few years16,18,35,84,137,138, it has become clear that X-ray spectromicroscopy uniquely addresses these issues.

For example, in polyurethanes containing common modifier particles, NEXAFS microscopy was able to differentiate two different modifier particles, provide an analysis of the uniformity of their chemical composition and to study their size distribution139. The materials studied were a polyurethane rich, poly-isocyanate poly-addition (PIPA)-based and styrene-acrylonitrile (SAN)-based particles dispersed in a TDI based, polyether rich polyurethane matrix. TEM of these materials shows both particles and does not discriminate between them. Figure 26 shows NEXAFS micrographs of this material at 285 eV (Fig. 26 A) and 287 eV (Fig. 26 B). The image at 285 eV shows both types of particles as dark dispersions, while at 287 eV only the SAN particles are highly absorbing relative to the matrix. The possibility of NEXAFS microscopy to image at different energies and change the relative contrasts allows for the

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discrimination between the two types of particles. A similar relative discrimination between the two types of particles has been achieved at the nitrogen absorption edge and, to a lesser extent, the oxygen edge. The chemical composition of the largest SAN particles has been determined and found to be uniform from particle to particle within the accuracy of the method, i.e. about 5%. It was determined that the SAN particles contain about a 20% polyether component.

The hard-segment urea distribution in polyurethane materials is strongly influenced by the relative amounts of urea and carbamate (urethane) present, which in turn is closely linked to formulation and processing. The urea/urethane ratio is thus an important quantity. In certain formulations, macro-phase separation occurs, but the composition of the precipitates formed has generally not been characterized on the submicron scale. Using polymer reference standards with known chemical composition, one can isolate the spectral signal corresponding to the urea and urethane components (typically only ~20% in systems of interest) and quantify the amounts of these two components through spectral fitting34,84 (see discussion in section 2.3). Based on similar quantitation, the composition and sharpness of precipitates in a high water TDI polyurethane foam has been assessed140. It was found that the precipitates were highly enriched in urea content (see Fig. 27) and had a fuzzy interface to the polyether-rich matrix. These observations provide a means to discriminate between various hypothesis on how these precipitates form

4.2.3. Multilayers Multilayer structures are used extensively in photographic applications, such as optical and x-ray photography and xerographic processes. For example, STXM studies of a test photoconductive thin film structure (courtesy Xerox Research Center of Canada) were used to investigate the degree of spatial uniformity of a N-containing charge transfer compound present in a protective polycarbonate (PC) capping layer above the image sensing layer. Figure 28 presents an image at 407 eV of the film, a N 1s linescan, and an analysis of the film uniformity. This single 5 minute measurement showed that the (~10%) N-content of the polycarbonate layer was distributed throughout the PC layer with a relatively uniform (±20%) distribution, but with some gradient in density towards the exterior of the thin film. By comparison, it was barely possible to detect any N, and thus impossible to analyze the spatial distribution of the charge transfer component, by either energy loss or X-ray fluorescence spectroscopy in a state-of-the-art JEOL 2010 analytical TEM. NEXAFS microscopy has also been useful in the characterization of several other multilayer polymer films used in photography. In one structure of four layers of 0.7 to 3 µm in thickness coated on a base layer, the images acquired clearly showed each of the layers and NEXAFS spectra of closely spaced areas indicated that there was not a significant interpenetration between two layers of interest, poly(styrene acrylonitrile) and a porous carbon black141,142. A second laminate consisting of nine layers on a base layer is

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shown in Fig. 29 142,143. While images obtained at four photon energies clearly show each of these layers, a microstructure of undetermined origin was found dispersed throughout the fourth layer at two photon energies. Domains are elongated parallel to the layer boundary. This microstructure was completely absent in the TEM, and is likely due to a compositional micro-phase separation or a preferential orientation of the aromatic groups present. Of particular interest in this line of work is the influence of coating conditions and thermal history upon the level of diffusion of small molecules and the level of interpenetration of adjacent polymer layers. 4.2.4. Elastomer Composites

Blends and composites based on Styrene-butadiene rubber (SBR) and butadiene rubber (BR) are widely used throughout the rubber industry in order to get a balance of properties which cannot be achieved through the use of a single elastomer. The mechanical properties of a rubber blend are determined, in part, by the phase morphology of the blend, which is typically formed using immiscible elastomers. The vast majority of these blends are heterogeneous on some length scale 144,145. These blends are further complicated by the addition of both fillers and curatives. In addition to determining the polymer phase structure itself, it is desirable to determine the distribution of fillers and curatives in each of the elastomer phases.

SBR and BR mixtures are used in most automobile and truck tires today. These application requirements are severe and tires are constantly being improved. For example, poly(isobutylene-co-4-bromomethylstyrene) can be blended with highly unsaturated general-purpose rubbers to impart unique barrier or dynamic properties and enhanced oxidative stability. The final properties of such a blend are the result of a complex series of compounding, mixing and curing stages. These stages profoundly impact the homogeneity of the mixed components, which include the polymers, the filler and the curatives. The characterization of the phase morphology of blends of commercial rubbers like polybutadiene, polyisoprene, poly(butadiene-co-styrene) and brominated poly(isobutylene-co-4-methylstyrene) continues to represent a problem for conventional techniques, especially in the case of filled rubbers. J. Dias et al. have thus started to employ NEXAFS microscopy for these applications and found that the phase morphology of various blends and the distribution of silica and carbon black fillers can be determined through the spectroscopic contrast inherent to these components146. A detailed characterization is essential to understanding and controlling the final properties of a heterogeneous elastomer blend.

4.3. Fibers

NEXAFS microscopy is also well suited to study the composition and structure of fibers. Ade and Hsiao have investigated the qualitative orientation of functional groups in poly(p-phenylene terephthalamide) (Kevlar®) fibers and found that aromatic groups are, on average, pointing radially outwards147.

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Smith and Ade subsequently characterized the relative quantitative orientational order in three different grades of Kevlar® fibers. Kevlar® 149 was found to be 2.3 and 1.6 times as radially orientated as Kevlar® 29 and Kevlar® 49, respectively148. This variation in orientation is relatively large, given that the crystallinity of all fiber grades is above 85%. Fig. 30 shows the spectra of Kevlar® 149 and 49 fibers, respectively, acquired from sample locations that had the electric field vector parallel or perpendicular to a radial position vector. These spectra clearly show that the dichroic ratio is smaller for the Kevlar® 49 fiber grade than for Kevlar® 149. It also nicely illustrates the sign reversal of the dichroism for the π* and σ* orbitals. More recent work on these fibers combined absorption spectroscopy at the carbon, nitrogen and oxygen edges. This combination, in conjunction with theoretical calculations, can estimate the absolute degree of radial orientation in these fibers. Preliminary analysis indicated that Kevlar® 149 is about 50% radiallly orientated149,150.

Kikuma et al. have used NEXAFS microscopy to study the effect of heat treatment on polyacrylonitrile fibers151. In heat-treated fibers, a clear core-rim structure was observed in images at several photon energies. A decrease in the nitrile group concentration was measured in the core of these fibers when compared to the untreated fibers. 4.4. Organic Geochemistry Most of the applied and fundamental problems in fuel chemistry and organic geochemistry are related to determining the molecular structure of geochemical polymers and relating this information to the time dependent response of such systems when subjected to environmental stresses, such as temperature and pressure. The difficulty in obtaining relevant information lies in the intrinsic microheterogeneous nature of geochemical polymers. The use of NEXAFS microspectroscopy provides the only analytical method available to date to probe the molecular chemistry of solid phase organics at the required spatial resolution. Over the past few years, Cody et al. have been exploring the use of NEXAFS microscopy to address a number of long standing problems in both fuel chemistry and organic geochemistry152-157. Cody et al. have, for example, followed the evolution of the molecular structure of an important microscopic constituent in coals, sporopollenin155. Determining its molecular structure and how it evolves when subjected to temperature and pressure over time holds the promise of addressing fundamental questions related to the mechanism and timing of oil generation. NEXAFS images of sporopollenin surrounded by a matrix of coal were acquired from a number of samples. Deconvolution of the point NEXAFS spectra from the sporopollenin in various rank coals allowed the tracking of the chemical structural evolution across a range of samples that had been subjected to progressively higher degrees of thermal metamorphism155. Cody et al. also studied the chemical structural evolution of various coals, Type III kerogens (vitrinite) ranging from low to high maturity, as they were subjected to blue

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light in an atmospheric environment156. Significant changes in the NEXAFS spectra could be observed and correlated to changes in the luminescence behavior of the coals. Traditionally, luminescence spectroscopy has been used to assess these changes, but that is a far more indirect measure of the chemical structural changes. With NEXAFS microscopy, the evolution of aromatic, carboxylic acid and other functional groups could be followed directly. The exposure to blue light in air resulted in the oxidation of the coal, with the dominant reaction being the formation of carboxyl groups. The photochemical oxidation pathways and kinetics varied significantly between the three maturity levels investigated.

Although high quality metallurgical coke is crucial to the fabrication of steel and a detailed understanding of the genesis of coke is highly desirable, the physical and chemical changes that occur during coking are difficult to characterize due to the fact that the critical stages involve nucleation at a very fine scale. Linear dichroism microscopy was thus utilized by Cody et al. to characterize the transformations that occur within coal during coking. In Fig. 31, we present the in-situ analysis of the chemistry and orientation of nematic phases in a quenched coke. In Figure 31A, a pair of C-NEXAFS spectra highlight the high degree of molecular orientation observed in these nematic phases. Primarily orientational and not compositional differences are responsible for the observed intensity differences in these spectra. The linear dichroism micrograph in Fig. 31B reveals a spectacular “Schlieren”-type image in one of these samples, where the contrast is based almost entirely on molecular orientation relative to the polarization of the x-ray beam. NEXAFS microscopy can also provide insight into the complexity of solid phase biomolecular polymers. High resolution images of 40 million-year-old wood are presented in Fig. 32. At a photon energy of 285.1 eV (Fig. 32A) the contrast is based on variations in the concentration of C=C bonded carbon, e.g. aromatic or olefinic carbon. Fig. 32B is acquired at a photon energy of 289.1 eV and contrast is primarily based on the concentration of carbon σ bonded with oxygen. NEXAFS microscopy reveals in enormous detail the high degree of chemical differentiation within the cell wall of wood. The composition of the primary and the two secondary cell walls, as well as the middle lamellae, could be assessed in detail from NEXAFS point spectra. Comparison to other wood samples that range in age from the present to 50 million years provides insight into the metamorphosis of carbohydrates over geological time scales157. 4.5. Meteoritics and Interplanetary Dust Particles

Interplanetary dust particles (IDPs) may have been an important source of prebiotic organic compounds necessary for the development of life on earth. IDPs from the Earth’s stratosphere are carbon-rich, containing on average 10-12 wt% carbon. However, the ratio of elemental to organic carbon is not well established. Flynn et al. have thus started to systematically map the spatial distribution of carbon, and to determine the carbon bonding states in ultra-

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microtomed thin sections of IDPs with NEXAFS microscopy158. All but one of 12 IDPs examined have significant quantities of organic carbon that occur in at least three morphologies159. An example of the “organic” signature detected is presented in Fig. 33.

Flynn et al. have also studied organic compounds associated with carbonate globules and rims in meteorites160. Specifically, ultrathin sections of

the ALH84001 meteorite were examined. For this meteorite, McKay et al.161 reported that polycyclic aromatic hydrocarbons (PCHs) were frequently found in the highest concentration in the regions rich in carbonates. PCHs are often produced by the decay of living material, and hence these finding could have a substantial impact on our identification of extra-terrestrial life. The association reported by McKay et al. has been produced with Laser Desorption Laser Ionization Mass-Spectrometry (LD-LIMS) with a spatial resolution of about 50 µm. Use of NEXAFS microscopy allows thus to examine the association of carbon bearing compounds with the mineral phases of ALH84001 with a spatial resolution about 1000 times better than LD-LIMS. Flynn et al. observed that there appear to be two different carbon-bearing phases, one of which is rather inhomogeneous. Correlation of the STXM carbon map with TEM mineralogy indicates that carbon rich phases are associated with fine grained magnetite and sulfide, and may also occur as veins or inclusions in feldspatic glass. 4.6. Soils and Environmental Studies Several research groups are in the early stages of utilizing NEXAFS microscopy to investigate the inhomogeneity of the organic compounds in soils and how metals, particularly heavy metal contaminants and other materials interact with different organic constituents of the soil and with microbes. The biochemical alteration of plant and animal residue in soils and aquatic systems results in the formation of natural organic polymers with a variety of molecular sizes and functional groups. Their concentration ranges from <1 ppm to as high as 40%, and consists of small chain molecules (e.g. acetate, citrate), organic macromolecules (e.g. proteins), and long chain, polyfunctional humic substances162.

The spectroscopic differences of isolated humic substances and the chemical micro-heterogeneity of the organic components in soil can be appreciated by the micrographs in Fig. 34. The sharp features near 288.5 and 290.5 eV and the doublet at 300 eV correspond to carboxylic and carbonate functionalities and the potassium L-edge, respectively. A variety of additional studies in this field are presently undertaken with the Stony Brook and ALS BL7.0 STXMs by research groups from SUNY@StonyBrook163, University of Central Florida164, and Göttingen165,166. 4.7. Other Applications

NEXAFS microscopy is also capable of characterizing biopolymers and biological systems. Zhang et al. first showed that NEXAFS microscopy can be

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used to discriminate the DNA and protein components of Chinese Hamster Ovary cells37, and subsequently mapped and measured the relative abundance of DNA and specific proteins (protamine 1 and protamine 2) in various mammalian sperm heads98. The data obtained for dried bull, stallion, hamster and mouse sperm suggests that the total nuclear protein to DNA ratio is similar in the sperm of many eutherian mammals. The total protamine content of sperm chromatin was constant among mammalian species, independent of the extent of expression of the protamine II gene. Zhang et al. concluded that protamine II replaces protamine I, rather than somehow binding to the protamine I complex. Based on experiments with human sperm, that NEXAFS microscopy may also prove useful for identifying protamine deficiencies in the sperm of infertile patients. Buckley et al. used NEXAFS microscopy to map protein and inorganic components in embedded bones101,167. A comparison between sections of ovariectomised and normal mouse bone showed significant differences in the density of protein between these two samples. In addition, a difference in micro-porosity could be inferred. These findings challenge the prevailing view that the matrix of osteoporotic-like bone always shows a normal composition. NEXAFS microscopy has also been used to characterize hydrated polymers. Mitchell et al. used a wet cell to visualize the spatial variation in crosslink density in thin sections from superabsorbent polymer (SAP) beads168. The swelling of SAP beads (polyacylic acid, sodium salt) and particularly of the more highly crosslinked surface shell depends on the cross-link density of the polymer. SAP beads with a high cross-link density will expand less in water and therefore will have a higher carbon density and image contrast. NEXAFS microscopy was used to visualize the core/shell structure in SAP beads prepared by chemically cross-linking the surface of SAP beads. Beads with different surface treatment have been compared. 5. Summary and Future Outlook NEXAFS spectroscopy at high spatial resolution is a powerful research tool that has already been used for a variety for scientific studies. While NEXAFS microscopy has moved well beyond proof-of-principle experiments and its capabilities are continually improving, the field has not yet reached any fundamental limits in performance and additional technological developments can be anticipated. For example, the far-field wavelength limited spatial resolution for zone plate based microscopes is about 3 nm for carbon K-edge energies, about an order of magnitude better than the resolution presently achieved. While significant technical challenges have to be overcome to achieve that kind of spatial resolution, we would like to point out that it is not the physics that prevents the field to go to higher spatial resolution. Similarly, several aberration corrected PEEMs with a projected spatial resolution of about 2 nm are planned or under construction169,170, while the best PEEMs to date have a spatial resolution of 20 nm. The technical challenges for the next generation PEEMs are also formidable. Several new operating modalities such

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as tomography171, and cryo-microscopy172 have been implemented in transmission which should prove very beneficial for future application. While there are still plenty of scientific questions that can be addressed with the present performance of existing NEXAFS microscopes, improved capabilities will further accelerate applications in polymer science and related fields. However, increased radiation dose as the spatial resolution improves is a serious issue for the study of polymeric materials. Procedural improvements or cryo-microscopy172 might have to be used extensively to minimize damage.

Additional applications might be made possible by the use of high energy resolution that reveals subtle spectral features. Studies on deuterated and hydrogenated PS showed, for example, that vibrational structure in these polymers exists and can be distinguished19. This might serve as an additional contrast mechanism in several applications. The establishment of a NEXAFS spectroscopy database similar to the spectroscopy database available in the IR community would greatly facility quantitation and analysis of materials for which prior knowledge is very limited. In parallel to the development of a spectroscopic database, improved theoretical modeling would greatly improve the ability to interpret NEXAFS spectra and the chemical structural information that they provide.

Based on the progress made to date and the improvements that can still be achieved, we are confidently looking forward to continuing growth of the NEXAFS spectroscopy and microscopy community and to exciting applications that have not been conceived yet. Acknowledgements We are grateful to S. Anders, G. Cody, A. Cossy-Favre, G. Flynn, A. P. Hitchcock, C. Jacobsen, S. Myneni, W. Meyer-Ilse, D. A. Winesett, and C. Zimba for sharing their results and thoughts, for permission to publishing their figures, as well as for providing some of the figures. Much of the work presented here is based on the work of a large group of people. We would like to thank them all, although we can not mention everybody by name. We would like to explicitly thank those that we have had closest contact with during the last few years, including S. Anders, G. Cody, T. Coffey, A. P. Hitchcock, C. Jacobsen, J. Kirz, G. Meggs, G. Mitchell, D. Pierson, M. Rafailovich, E. Rightor, A.P. Smith, J. Sokolov, R. Spontak, T. Warwick, S. Wirick, D. A. Winesett, and C. Zimba. They have made our work very pleasurable and exciting. We attempted to present the wide range of applications of NEXAFS spectroscopy and microscopy of organic materials and have drawn heavily on our own results. The presented material is not intended to cover or reference all the work being performed. We regret if we should have missed some interesting and important work.

With few exceptions, the data presented has been acquired with either the SUNY-SB STXM at the National Synchrotron Light Source, the B7.0 STXM or the PEEM at the Advanced Light Source. The NSLS and ALS are supported by

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the Department of Energy. Some of the work presented and the writing of this manuscript is supported in part by a NSF National Young Investigator Award (DMR-9458060) and Dow Chemical. Tables Table 1: mol% of FORMULA UNITS of indicated component

Species

polyether Urea carbamate

Predicted

Fit Predicted

Fit Predicted

Fit

258 86.7 86 11.8 12 1.6 2 259 90.6 89 6.3 8 3.1 3 260 89.9 90 3.6 4 6.5 6

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Figure Captions Figure 1. Electronic schematic for the NEXAFS photoabsorption spectrum of poly(styrene-r-acrylonitrile). The C 1s binding energies were obtained from XPS173 and orbital energies of the C 1s → LUMO(π*) transition of the phenyl (A) and acrylonitrile (B) functional groups were obtained from ab initio calculations174. Figure 2. C 1s NEXAFS spectra of poly(ethylene terephthalate) (PET), recorded by C 1s Electron Energy Loss Spectroscopy (EELS), recorded using a LaB6 electron source17 and a field emission gun (VG-HB501)16, in comparison to the NEXAFS spectrum recorded by total electron yield16 and transmission18. Figure 3 C 1s NEXAFS spectra for a series of unsaturated polymers, recorded in transmission in the SUNY-SB STXM microscope at the National Synchrotron Light Source. Figure 4 C 1s NEXAFS spectra of a series of saturated polymers, recorded in transmission in the SUNY-SB STXM microscope at the National Synchrotron Light Source. Figure 5 C 1s NEXAFS spectrum of (A) poly(ethylene terephthalate) (PET)18 in comparison to the C 1s ISEELS spectra of (B) dimethyl terephthalate18, (C) ethyl benzoate 32 and a simulation formed by the sum of the spectrum of benzene and formic acid (D)32. Figure 6 C 1s NEXAFS spectra of three polyurethane model polymers (denoted B0, B2, BM) that differ in their hard segment/soft segment composition, in comparison to the C 1s ISEELS spectrum of ethyl N-phenyl urethane and the C 1s NEXAFS spectrum of 2-propane diol34. Figure 7 C 1s NEXAFS spectrum of poly(propylene oxide) and 2-propane diol (solid phase spectrum), in comparison to the C 1s ISEELS gas phase spectrum of 2-propane diol 34. Figure 8 C 1s NEXAFS spectrum of Glysine and Tyrosine, as well as the C 1s NEXAFS spectrum of the Gly-Tyr dimer overlaid with the simulated spectrum on the dimer produced by addition of the Glysine and Tyrosine NEXAFS spectra. (Figure courtesy of C. Jacobsen, SUNY@Stony Brook.) Figure 9 Polarization-dependent C 1s NEXAFS spectra of highly oriented polytetrafluoroethylene from Ziegler et al.58 Spectra are recorded at angles of

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α = 0°, 22.5°, 45°, 67.5° and 90° between the electric field vector (eE) and the chain axis. (Reproduced from ref. 58.) Figure 10 C 1s NEXAFS spectrum of a blend of polycarbonate (PC) and poly(butylene terephthalate) (PBT), in comparison to the C 1s NEXAFS spectra of PC and PBT homopolymer. (Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 11 (left) C 1s NEXAFS spectra of TDI polyurethane, TDI polyurea and poly(propylene oxide), selected to represent the spectra of the urethane, urea and polypropylene oxide linkages present in polyurethane foams. (right) Quantitative analysis of C 1s NEXAFS spectra of 3 different model polyurethane polymers of varying urea and urethane content (denoted 258, 259, and 260. See text and table 1 for details.). The polyurethane data is indicated by filled circle symbols, the best fit by the thicker solid line and the amounts of the individual polyether (E), TDI-polyurea (U), and TDI-polyurethane (C) component spectra required to construct that best fit are indicated by the thin solid lines. (Figure reproduced from ref. 84) Figure 12 Schematic of a Scanning Transmission X-ray Microscope (STXM) and a conventional Transmission X-ray Microscope (TXM). (Figure courtesy of W. Meyer-Ilse, CXRO.) Figure 13 Schematic of a Photoemission Electron Microscope (PEEM). (Figure courtesy of S. Anders, ALS, and B. Tonner, Uni. Central Florida.) Figure 14 Reference spectra of PS and PMMA and transmission images of PS/PMMA sample at the four energies indicated. Although the absorption in some x-ray micrographs is dominated by one or the other chemical substance, e.g. 285 eV is dominated by PS absorption, each images typically has a weighted absorption, based on the cross sections at the particular photon energy and the abundance of the substance. (Figure courtesy of D.A. Winesett, NCSU. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 15 Raw transmission images are normalized to optical density (OD) images, which are now linear in the absorption coefficients. A set of OD images can subsequently be processed via a singular value decomposition (SVD) procedure to yield quantitative composition maps of the known constituents. (Figure courtesy of D.A. Winesett, NCSU. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 16 PEEM images of a stylus track on a polyimide surface at two different orientations at a photon energy of 285.3 eV. Contrast is based on orientational and alignment differences of the polymer chains inside and outside

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the track. The uniform gray outside the track corresponds to isotropic chains104. (Figure courtesy of A. Cossey-Favre.) Figure 17 Dichroic NEXAFS spectra of the stylus track in Fig. 16. Spectral differences between the spectra are due to differences in orientation and alignment of the polymer chains inside and outside the track104. (Figure courtesy of A. Cossey-Favre.) Figure 18 Wear tracks in the carbon coating of a ultra-smooth surface of a hard disk, and NEXAFS spectra inside and outside the wear tracks revealing very different compositions at these locations. (Figure courtesy of S. Anders, ALS.) Figure 19 Images and spectra of a trilayer structure consisting of a 60 nm film of polystyrene (PS) coated on each side with 30 nm of SiOx. Annealing results in the observed morphology. Profiles and spectra show that significant amounts of PS remain in the thin regions. (Figure courtesy of A. P. Hitchcock, McMaster University. Data acquired with the BL7.0 STXM microscope at the Advanced Light Source.) Figure 20 PS mass thickness (left column), PMMA mass thickness (middle column) and total thickness maps (right column) of a nominally 143 nm thick 50/50 w/w% PS/PMMA blend annealed for (A-C) 0 min, (D-F) 2 min, (G-I) 10 min, (J-L) 30 min, (M-O) 2 hrs, and (P-R) 1 week. All images are individually scaled for good contrast, with Black = 0 and White = maximum thickness. The maximum thickness of the films increases from 145 nm in Fig. 20A, to 460 nm in Fig. 20R. (Figure reproduced from ref. 100) Figure 21 PS (a), PBrS (b), and total thickness (c) maps of dewetting PBrS/PS bilayer after one week of annealing at 180°. Line profiles below the images are from the location indicated by horizontal lines in the images. (Figure reproduced from ref. 99.) Figure 22 PEEM images and spectra from the same sample shown in Fig. 21. The combination of the STXM and PEEM data allowed to deduce the three dimensional morphology with certainty. (Adopted from ref. 99) Figure 24 (A) TEM image of PS/PHIC diblock copolymer and small area electron diffraction pattern averaged over several lamellae. (B) NEXAFS image at 285 eV (PS dark), (C) NEXAFS image at 288.5 eV, where the PHIC layers have altering contrast due to differences in orientation. (Figure courtesy of C. Zimba, NIST and E. Thomas, MIT. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 25 STXM images acquired at 286.7 eV of PET/VectraTM blends

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differing in composition (in w/w PET/VectraTM): (A) 75/25, (B) 90/10 and (C) 99/1. High image contrast allowed easy assessment of the size distribution of VectraTM dispersions within the PET matrix. The internal structure of the VectraTM domains is due to molecular orientation as well as thickness variations (see Fig. 26). (Figure reproduced from ref. 133.) Figure 26 STXM images acquired at 286.7 eV of a 99/1 w/w PET/VectraTM blend subjected to post-milling melt pressing. Images (A) and (B) have been converted to optical density. In images (A) and (B), the electric polarization vector (E ) is rotated by 90° with respect to each other, as indicated. Changes in the relative intensity in these images are primarily due to anisotropic molecular orientation. The ratio of these images (C) reveals the linear dichroism of the specimen. Small VectraTM domains appear gray and possess no discernible orientation, whereas the large dispersion exhibits a measurable degree of molecular orientation (black and white areas) due to the nematic nature of this liquid crystalline polymer. (Figure reproduced from ref. 133.) Figure 26 C 1s NEXAFS micrographs of polyurethane polymer with two different types of filler particles; PIPA and SAN (see text). (a) image at 285.3 eV showing both types of filler particles, (b) image at 287 eV, showing only the SAN particles with high contrast, while the contrast between the PIPA and the matrix almost vanished. (Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 27 Quantitative determination of the composition of the matrix and precipitates in a water-rich TDI polyurethane foam based on fitted point spectra. (Figure courtesy of A. P. Hitchcock, McMaster University. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 28 (left) STXM image of a photoconductive thin film structure at 407 eV. From bottom to top the layers are: a poly(ethylene-terephthalate) base, a conductor, a pigment layer, a polycarbonate (PC) cap with additive, and an epoxy. (center, lower) linescan - optical density as a function of energy and position along the indicated line. (upper) total and N 1s OD at 407 eV, obtained by integrating signals in the indicated regions. (Figure courtesy of A. P. Hitchcock, McMaster University. Data acquired with the BL7.0 STXM microscope at the Advanced Light Source.) Figure 29 C- NEXAFS micrograph at a photon energy of 295.2 eV of part of a polymer laminate with unexpected microstructure in layer #4. (Figure courtesy of C. Zimba, NIST. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.)

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Figure 30 C 1s NEXAFS spectra of transverse sections of Kevlar® 149 and Kevlar® 49 fibers at the two locations indicated. The difference between the two spectra for each fiber grade provides the dichroic ratio, which is clearly smaller for Kevlar® 49 than for Kevlar® 149. (Figure reproduced from ref. 148.) Figure 31 NEXAFS spectroscopy (A) and dichroism microscopy (B) of a coke. Spectral differences and image contrast are almost exclusive due to variations in orientation and not chemical composition. (Figure courtesy of G. Cody, Geophysical Laboratory, Carnegie Institution of Washington. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 32 NEXAFS images of 40 million year old wood at the photon energies indicated. The variation of aromatic-olefinic concentration (A) as well as the carbohydrate concentration (B) can be mapped157. (Figure courtesy of G. Cody, Geophysical Laboratory, Carnegie Institution of Washington. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.)

Figure 33 Carbon 1s NEXAFS spectrum from the rim of a meteorite, showing spectral features associated with organic carbon, rather than atomic carbon, in the sample. (Figure courtesy of G. Flynn, SUNY@StonyBrook-Plattxburgh. Data acquired with the SUNY-SB STXM microscope at the National Synchrotron Light Source.) Figure 34 C-NEXAFS spectra of the isolated (left) and pristine soil humic substances (right). Center image is from pine ultisol aggregates in water. This soil sample contains Fe-oxides and clays, and the organic-C concentration is about 4.5 %. The soil solution has a pH of 5.0. The NEXAFS spectra shown on the right are for different locations in the soil aggregate. The sharp peaks at ~ 288.5, 290.5, and the doublet at 300 eV correspond to the carboxylic, carbonate and the potassium L2 & L3 edges, respectively. The low-energy shoulder at 287 eV corresponds to the ketonic, and aliphatic C-H groups. (Figure courtesy of S. Myneni, LBNL. Data acquired with the BL7.0 STXM microscope at the Advanced Light Source.)

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118. S. Walheim, M. Boltau, J. Mlynek, G. Krausch, and U. Steiner, "Structure Formation via Polymer Demixing in Spin-Cast Films", Macromol 30, 4995 (1997).

119. A. Karim, T. M. Slawecki, S. K. Kumar, J. F. Douglas, S. Satija, C. C. Han, T. P. Russell, Y. Liu, R. Overney, J. Sokolov, and M. H. Rafailovich, "Phase-Separation Induced Surface Patterns in Thin Polymer Blend Films", Macromol. 31, 857 (1998).

120. P. Müller-Buschbaum, S. A. O'Neill, S. Affrossman, and M. Stamm, "Phase Separation and Dewetting of Weakly Incompatible Polymer Blend Films", Macromol. 31, 5003 (1998).

121. A. Cossy-Favre, J. Diaz, S. Anders, H. Padmore, Y. Liu, M. Samant, J. Stohr, H. Brown, and T. P. Russell, "Photoelectron emission microscopy and its application to the study of polymer surfaces", Acta Physica Polonica A. 91, 923 (1997).

122. D. Slep, J. Asselta, M. H. Rafailovich, J. Sokolov, D. A. Winesett, A. P. Smith, H. Ade, Y. Strzhemechny, S. A. Schwarz, and B. B. Sauer, "Phase Separation of Polystyrene and Bromo-polystyrene Mixtures in Equilibrium Structures in Thin Films", Langmuir 14, 4860 (1998).

46

123. D. A. Winesett, H. Ade, J. Sokolov, M. Rafailovich, and S. Zhu, "Substrate Dependence of Morphology in Thin Film Polymer Blends", Polymer International (in press) (2000).

124. A. P. Smith, H. Ade, D. Slep, S. Qu, M. Rafailovich, J. Sokolov, G. Halada, S. A. Schwarz, and Y. Strzhemechny, "Investigation of Spinodal Decomposition in Polymer Thin Films at Surfaces using X-ray Microscopy", MRS 1996 Fall Meeting Abstracts (1996).

125. A. P. Smith, H. Ade, D. Slep, S. Qu, M. Rafailovich, and J. Sokolov, "Investigation of the Spinodal Decomposition of Constrained Thin Polymer Films with X-ray Microscopy", NSLS Activity Report for 1996, Brookhaven National Laboratory (1997).

126. A. J. Wagner and J. M. Yeomans, "Breakdown of Scale Invariance in the Coarsening of Phase-Separating Binary Fluids", Phys. Rev. Lett. 80, 1429 (1998).

127. S. Zhu, Y. Liu, M. H. Rafailovich, J. Sokolov, D. Gersappe, D. A. Winesett, and H. Ade, "Confinement Induced Miscibility in Polymer Blends", Nature 400, 49 (1999).

128. D. Slep, J. Asselta, M. H. Rafailovich, J. Sokolov, D. A. Winesett, A. P. Smith, H. Ade, and S. Anders, "The Effect of an Interactive Surface on the Equilibrium Contact Angle in Bilayer Polymer Blends", Langmuir (in press ) (2000).

129. C. Zimba, E. Thomas, and C. Ober, (in prepartion) (1999). 130. J. T. Chen, E. L. Thomas, C. K. Ober, and G. Mao, Science 273, 343 (1996). 131. H. Ade, A. Smith, S. Cameron, R. Cieslinski, C. Costello, B. Hsiao, G. Mitchell, and

E. Rightor, "X-Ray microscopy in polymer science: Prospects of a "new" imaging technique", Polymer 36, 1843 (1995).

132. H. Ade, A. P. Smith, G. R. Zhuang, B. Wood, I. Plotzker, E. Rightor, D.-J. Liu, S.-C. Lui, and C. Sloop, "X-ray Microscopy of Multiphase Polymeric Materials", Mater. Res. Soc. Symp. Proc. 437, 99 (1996).

133. A. P. Smith, C. Bai, H. Ade, R. J. Spontak, C. M. Balik, and C. C. Koch, "X-ray Microscopy of Novel Thermoplastic/Liquid Crystalline Polymer Blends by Mechanical Alloying", Macromol. Rapid Commun. 19, 557 (1998).

134. A. P. Smith, H. Ade, R. J. Spontak, and C. C. Koch, "Morphological and Chemical Characterization of a Mechanically Alloyed Rubber Toughened PMMA with X-ray Spectroscopy", Microscopy and Microanalysis 4 S-2, 142 (1998).

135. A. P. Smith, J. H. Laurer, H. W. Ade, S. D. Smith, A. Ashraf, and R. Spontak, "X-ray Microscopy and NEXAFS Spectroscopy of Macrophase-Separated Random Block Copolymer/Homopolymer Blend", Macromol. 30, 663 (1997).

136. A. P. Hitchcock, S. G. Urquhart, H. Ade, E. G. Rightor, and W. Lidy, "Quantitative Chemical Speciation of Multi-Phase Polymers Using Zone Plate X-ray Microscopy", Microscopy and Microanalysis 4 S-2, 808 (1998).

137. S. G. Urquhart, A. P. Hitchcock, R. D. Leapman, R. D. Priester, and E. G. Rightor, "Analysis of Polyurethanes Using Core Excitation Spectroscopy. Part I: Model Polyurethane Foam Polymers", J. Polym. Sci. B, Polym. Phys 33, 1593 (1995).

138. S. G. Urquhart, A. P. Hitchcock, E. G. Rightor, H. Ade, and A. P. Smith, "Chemical Speciation by NEXAFS Spectromicroscopy: Insights from Molecular Modelling of Polymers", Mater. Res. Soc. Symp. Proc. 437, 243 (1996).

139. A. P. Hitchcock, T. Tyliszczak, E. G. Rightor, G. E. Mitchell, M. T. Dineen, W. Lidy, R. D. Priester, S. G. Urquhart, A. P. Smith, and H. Ade, "Filler Particles in Polyurethanes Studied by Soft X-ray Spectromicroscopy", (to be submitted) (1999).

140. E. G. Rightor, S. Urquhart, A. Hitchcock, H. Ade, A. P. Smith, G. Mitchell, R. Priester, and W. Lidy, "Identification and quantitation of urea precipitates in flexible polyurethanes by NEXAFS microscopy", (to be submitted) (2000).

141. C. Zimba, A. P. Smith, and H. Ade, "Characterization of Polymer Interfaces Using X-ray Microscopy", (to be submitted) .

47

142. C. Zimba, A. P. Smith, and H. Ade, NSLS Activity Report for 1996, Brookhaven National Laboratory (1997).

143. C. Zimba, A. P. Smith, and H. Ade, "X-Ray Microscopy of Polymer Laminates: NEXAFS Microspectroscopy and Imaging", (to be submitted) .

144. C. M. Roland, Rubber Chem. Technol. 62, 456 (1989). 145. W. M. Hess, C. R. Herd, and P. C. Vegvari, Rubber Chem. Technol., 66, 329 (1993).

146. J. Dias, S. G. Urquhart, H. Ade, and P. Stevens, "Chemical Sensitive Imaging of Tire

Blends Using X-ray Microscopy", 155th Meeting of the Rubber Division, American Chemical Society, Chicago , paper no. 30 (1999).

147. H. Ade and B. Hsiao, "X-ray Linear Dichroism Microscopy", Science 262, 1427 (1993).

148. A. P. Smith and H. Ade, "Quantitative Orientational Analysis of a Polymeric Material (Kevlar Fiber) with X-ray Microspectroscopy", Appl. Phys. Lett. 69, 3833 (1996).

149. H. Ade, A. P. Smith, H. Garcia, T. Warwick, S. Cerasari, and B. Tonner, "Lateral Orientational Characterization of Kevlar Fibers with X-ray Linear Dichroism Microscopy", Bull. Am. Phys. Soc. 42, 47 (1997).

150. H. Ade, A. P. Smith, A. Garcia, and S. Urquhart, "Quantitative orientational analysis of polymers at high spatial resolution: The lateral orientational order of poly (p-phenylene terephthalamide) fibers", (in preparation) (2000).

151. J. Kikuma, T. Warwick, H. J. Shin, J. Zhang, and B. P. Tonner, "Chemical state analysis of heat-treated polyacrylonitrile fiber using soft X-ray spectromicroscopy", J. Electron. Spectrosc. Relat. Phenom. 94, 271 (1998).

152. G. D. Cody, R. E. Botto, H. Ade, S. Behal, M. Disko, and S. Wirick, "Inner shell spectroscopy and imaging of a sub bituminous coal: in situ analysis of organic and inorganic microstructure using C(1s)-, Ca(2p), and Cl(2s) NEXAFS", Energy & Fuels 9, 525 (1995).

153. G. D. Cody, R. E. Botto, H. Ade, S. Behal, M. Disko, and S. Wirick, "C-NEXAFS Microanalysis and Scanning X-ray Microscopy of Microheterogeneities in a High Volatile A Bituminous Coal", Energy & Fuels 9, 153 (1995).

154. R. E. Botto, G. D. Cody, J. Kirz, H. Ade, S. Behal, and M. Disko, "Selective Chemical Mapping of Coal Microheterogeneity by Scanning Transmission X-ray Microscopy", Energy & Fuels 8, 151 (1994).

155. G. D. Cody, R. E. Botto, H. Ade, and S. Wirick, "The application of soft x-ray microscopy to the in-situ analysis of sporinite in coal", Int. J. Coal Geol. 32, 69 (1996).

156. G. D. Cody, H. Ade, S. Wirick, G. D. Mitchell, and A. Davis, "Determination of chemical-structural changes in vitrinite accompanying luminescence alteration using C-NEXAFS analysis", Organic Geochemistry 28, 441 (1998).

157. G. D. Cody, H. Ade, K. Anderson, J. Brandes, and S. Wirick, "Probing Chemical Differentiation in the Cell Membrane of Recent and Ancient Wood (45 and 70 Ma) with Soft X-rays", (to be submitted) (2000).

158. G. J. Flynn, L. P. Keller, C. Jacobsen, and S. Wirick, "Carbon Mapping and Carbon-XANES Bonding State Measurements on Inter-Planetary Dust Particles", Lunar and Planetary Science XXIX, Abstract #1159 (1998).

159. G. J. Flynn, L. P. Keller, C. Jacobsen, and S. Wirick, "Organic Carbon Contributed by the Accretion of Interplanetary Dust Over the Earth's History", Amer. Geophys. Union Fall 1998 Meeting (1998).

160. G. J. Flynn, L. P. Keller, M. A. Midler, C. Jacobsen, and S. Wirick, "Organic Compounds Associated with Carbonate Globules and Rims in the ALH84001 Meteorite", Lunar and Planetary Science XXIX, Abstract #1156 (1998).

161. D. S. McKay, E. K. Gibson, Jr., K. L. Thomas-Keprta, H. R. Vali, C.S., S. J. Clemett, X. D. F. Chillier, C. R. Maechling, and R. N. Zare, "Search for past life on Mars:

48

possible relic biogenic activity in martian meteorite ALH84001", Science 273, 924 (1996).

162. F. J. Stevenson, "Humus Chemistry" (Wiley, New York, 1994). 163. U. Neuhäusler, S. Abend, C. Jacobsen, and G. Lagaly, "Soft X-ray spectromicroscopy

on solid-stabilized emulsions", Colloid Polym. Sci. 277, 719 (1999). 164. J. Rothe, E. M. Kneedler, K. Pecher, B. P. Tonner, K. H. Nealson, T. Grundl, W.

Meyer-Ilse, and T. Warwick, "Spectromicroscopy of Mn distributions in micronodules produced by biomineralization", J. Synchrotron Radiation 6, 359 (1999).

165. T. Preis and J. Thieme, "Dynamical Studies of Aqueous Clay Mineral Dispersions by X-ray Microscopy", Langmuir 12, 1105 (1996).

166. J. Thieme, S. Abend, and G. Lagaly, "Aggregation in Pickering emulsions", Colloid Polymer Science 77, 257 (1999).

167. C. J. Buckley, N. Khaleque, S. J. Bellamy, M. Robins, and X. Zhang, "Mapping the organic and inorganic components of tissue using NEXAFS", Journal de Physique IV 7 (C2) part 1, 83 (1998).

168. G. E. Mitchell, S. G. Urquhart, L. Wilson, M. Dineen, E. G. Rightor, A. P. Hitchcock, U. Neuhaeusler, H. W. Ade, W. Meyer-Ilse, J. T. Brown, and T. Warwick, "Analysis of the spatial variation of crosslink density in superabsorbent polymers", Advanced Light Source Compendium 1997, LBNL Report 41658 (1998).

169. R. Wichtendahl, R. Fink, H. Kuhlenbeck, D. Preikszas, H. Rose, R. Spehr, P. Hartel, W. Engel, R. Schlogl, H. Freund, A. Bradshaw, G. Lilienkamp, T. Schmidt, E. Bauer, G. Benner, and E. Umbach, "SMART: An aberration-corrected XPEEM/LEEM with energy filter", Surf. Rev. Lett. 5, 1249 (1998).

170. R. Fink, M. R. Weiss, E. Umbach, D. Preikszas, H. Rose, R. Spehr, P. Hartel, W. Engel, R. Dogenhardt, R. Wichtendahl, H. Kuhlenbeck, W. Erlebach, K. Ihmann, R. Schlögl, H.-J. Freund, A. M. Bradshaw, G. Lilienkamp, T. Schmidt, E. Bauer, and G. Benner, "SMART: a planned ultrahigh-resolution spectromicroscope for BESSY II", J. Electron Spectrosc. Relat. Phenom. 84, 231 (1997).

171. Y. Wang, C. Jacobsen, J. Maser, and A. Osanna, "Soft x-ray microscopy with a cryo STXM: II. Tomography.", J. Microscopy in press (2000).

172. J. Maser, A. Osanna, Y. Wang, R. Fliller, C. Jacobsen, J. Kirz, S. Spector, M. Weigel, B. Winn, and D. Tennant, "Soft x-ray microscopy with a cryo-STXM: Instrumentation, Imaging and spectroscopy", Journal of Microscopy (in press) (1999).

173. G. Beamson and D. Briggs, "High resolution XPS of organic polymers: the Scientia ESCA300 database" (John Wiley & Sons Ltd, 1992).

174. S. G. Urquhart, unpublished data (1999).

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Energy (eV)

285 290 295

CN

Bin

ding

Ene

rgy 284

285

286

287

Orb

ital E

nerg

y ( ε

)

-4

-3

-2

CN

A B

AB

Figure 1 Figure 2

Energy (eV)285 290 295

NEXAFS (STXM)NSLS X1A

NEXAFS (TEY)NSLS - U1A

TEM-EELSLaB6 JEOL 2000FX

TEM-EELSFE VG-HB501

CO

OC

O

O CH2 CH2

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Photon Energy (eV)285 290

CH3

CH3

O CO

O

CCO

OO

O (CH2)2

CC

O O

N

H

N

H

NCH

ONH

NNH

CO

O CH2C

H

OO

4

NCH

ONH

CH2

NC

H

O

O CH2 NH

C

O

O CH2 4

TDI urethane

TDI urea

MDI urea

MDI urethane

KevlarTM

PET

PC

OCO

C

O

Oy

x

VectraTM

Figure 3

Photon Energy (eV)285 290

CN

n m

Br

2278

Cl

SAN

PBrS

PS

NP

PBD

PI

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Energy (eV)

284 286 288 290 292

PEO

EPR

PP

PE

O

PIB

Figure 4

O PPO

Energy (eV)285 290

C CO

OO

O CH2CH2

C CO

OO

O

CO

O

CH2 CH3

HC

OH

O+ 2x

A: PET

B: 1,4 DMP

Figure 5

C:

D:

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Energy (eV)

285 290 295

NNCC

OO

O O H H

NC

O

O

H

NCOO

CN

O

H

OCN

~34

NNCC

OO

O O

H H

OHOH+

NC

O

O

H

NCOO

C

O

OCN

~34

BM Model MDI Polyurethane

B2 Model MDI Polyurethane

B0 Model MDI Polyurethane

OHOH

NC

O

O

H

B0

B2

BM

*0.8

*0.8

Figure 6

Energy (eV)

285 290 295

Osc

illat

or S

tren

gth

Per

Car

bon

Ato

m (

10-2

eV

-1)

0

1

2

3 O

OHOH

solid

OHOH

gas

Figure 7

Figure 8

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Figure 9

Energy (eV)

284 288 292 296

PBT

PC

PC/PBT

Opt

ical

Den

sity

Figure 10

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Energy (eV)285 290 295

NCH

ONH

NNH

CO

O CH2CH

OO 4

O

C

U

E

Figure 11

258

259

260

E

U

C

E

U C

E

U C

0

1

Osc

illat

or S

tren

gth

Per

Car

bon

Ato

m (

10-2

eV

-1)

0

1

Energy (eV)

285 290 295

0

1

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Figure 12

Sample

MagnifiedImage

X-ray Beam

Objective Lens Projective Lens

Screen

MCP

Aperture

Figure 13

Figure 14

Figure 16

(a) Line // E285.3 eV

field of view=76µµm

(b) Line ⊥⊥ E285.3 eV

ΦΦ″″ =171°

ΦΦ⊥⊥ ==31°

2 8 0 2 9 0 3 0 0 3 1 0 3 2 0

Abs

orpt

ion

coef

ficie

nt (

0

5

1 0

1 5

2 0

P M M A

P S

Pre C edge6.0µm

C 1saaπ∗π∗C=C C 1saaπ∗π∗C=O

Photon Energy (eV)

Abs

orpt

ion

coef

fici

ents

, α(µ

m-1

) Continuum

Transmission x-ray micrographs

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Transmission x-ray spectromicrographs

Optical density mapsOD = - ln(I/IO)

Sample areas, I

Open areas IO

PMMAPMMAPSPS

PMMAPMMAPSPS

PMMAPMMAPSPS

PMMAPMMAPSPS

ttOD

ttOD

ttOD

ttOD

⋅+⋅=

⋅+⋅=

⋅+⋅=

⋅+⋅=

0.3100.3100.310

1.2881.2881.288

2.2852.2852.285

8.2818.2818.281

αα

αα

αα

αα

Absorption coefficients αα from NEXAFS reference spectra

Solve by SVD for each component’s

thickness

Compositional maps

PMMA PS Total

Figure 15 Figure 17

1.5

1.0

0.5

0.0

330320310300290280Photon enery ( eV)

E parallel = I y

E perpendicular = I x

C K absorption edge

290288286284

285.3 eV

O

O

N

inset 2

inset 1285.3 eV

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

1000

1500

2000

2500

3000

280 285 290 295 300 305 310 315

energy (eV)

Line profile across stripes

Position (µµm)

0 5 10 15

Opt

ical

den

sity

(ar

b. u

nits

)

0.0

0.5

1.0

1.5

2.0 285 eV

281 eV

Spectra from Line Scans

Energy (eV)

284 288 292 296

Opt

ical

den

sity

(ar

b. u

nits

)

0.0

0.5

1.0

1.5

2.0

PS

"SiOx"

avg.

10 µµmA: 281 eV

B: 285 eV

Figure 19

Figure 18(a)

(b)

(c)

outside wear track

edge ofwear track

between rails

690 695 700 705 710energy (eV)

A B C

D E F

G A

H A

I A

J K L

M N O

P Q R 1 week

2 hrs

30 min30 min

2 min

0 min

6.0 µm

6.0 µm

4.0 µm

4.0 µm

2.0 µm

2.0 µm

G H I 10 min

Figure 20

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Thi

ckne

ss (

µ m)

0 . 0

0 . 1

0 . 2

0 . 3Max=299 nm Max=302 nmMax=232 nm

a b c

8.0µm

Figure 21

Figure 22

500 nm 1500 nm 1500 nm

A B C

Figure 23

A B C

3 µm 2 µm 2 µm

Figure 24

A B C

2 µmEE

Figure 25

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Figure 26

2 µm

A B

2 µmEnergy (eV)

284 286 288 290 292

Osc

. Str

. (10 urea

ether

urethane

Ether 40 (4)Urea 55 (5)Urethane 5 (3)

precipitate

PRECIPITATES : ΣΣ 4 spectra

Energy (eV)284 286 288 290 292

Osc

. Str

. (10

ether

urea

urethane

matrix

Ether 60 (5)Urea 30 (4)

Urethane 10 (3)

MATRIX : ΣΣ 4 spectra

Figure 27

Figure 28

4

5 µm

Position across PC layer

1 2 3 4 5 6 7

Inte

nsi

ty (

OD

) at

407

eV

0.0

0.1

0.2

0.3

0.4

0.5

0.6

N 1s & C 1s

N 1s only

1

2

3

5

67

E (eV)

Figure 29

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

Energy (eV)290 300 310

Opt

ical

Den

sity

0

1

2

3

4

5

Kevlar 149

Kevlar 49

r→

r→→→

rE ||(dot)

→→

⊥ rE(solid)

E

Figure 30

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

Polarization Parallel and Perpendicular

280 285 290 295 300 305 310

eV

5 µm E = 288.5 eVFigure 31

A B

4 µm E=285.5 eV E=289.5 eV

A B

Figure 32

Ade and Urquhart NEXAFS Spectroscopy and Microscopy of Natural and Synthetic Polymers

E n e r g y ( e V )

2 8 0 2 8 5 2 9 0 2 9 5 3 0 0

Opt

ical

Den

sity

0 . 0

0 . 1

0 . 2

0 . 3

0 . 4

0 . 5

Fig. 33

11

3

2

0

.5

1

1.5

2

280 290 300

Photon Energy ( eV )

fulvic acid

humic acid

Abs

orpt

ion

Inte

nsity

(arb

itrar

y un

its)

Photon energy (eV)

1

2

280 285 290 295 300 305

3

Fig. 34