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How sticky droplets behave when there's a lot of them

I've been writing my dissertation, so updates have been scarce, but I can report some interesting papers I've read along the way.

There is a great paper by Ryotaro Shimizu and Hajime Tanaka that explains a new mechanism for how immiscible droplets within a fluid merge into bigger droplets, and has universal applications from emulsions to cosmetics and foods.  Immiscible means two fluids cannot mix.  The models described here work are applicable for low viscosities.  This work is based on computer simulations, but the authors claim there are experiments to back up their claims.  The paper has been recently published in Nature Communications, so time will tell on what further peer review will reveal. 

Shimizu and Tanaka focus on phase separation, which is a fundamental physical phenomenon of multiple bodies collectively interacting.  Most familiar encounter I can think in everyday life is olive oil droplets merging into a single mass in a cooking pot.  The process is called coarsening, and the oil droplets can be thought of as domains in a continuous medium called water.  The way by which droplets coalesce is dependent on their initial concentration.  The concentration changes the mechanism by which they interact.  

When the concentration is low, droplets merge via a mechanism called Ostwald ripening.  This is as old as the hills (actually since the 1890s), and works by smaller droplets steadily diffusing into large droplets at the expense of the smaller droplets.  

Figure 1.  Small droplets spontaneously dissolve due to their instability.  The material then diffuses into larger droplets, which are more energetically stable.    Source: Wikipedia

Another mechanism is works when there are enough droplets to occupy about 50% of the total volume in a given object.  There, hydrodynamic forces coarsen the droplets into a so-called bicontinuous pattern, as seen in Figure 2.  Bicontinuous refers to two continuous phases.   You can see the black and white patterns continuing beyond the box in the moments of the animation.  The merging is so rapid that normal diffusion is dwarfed by the swift currents that result in the final structure of the animation. 

Figure 2.  Emergence of bicontinuous structure.  Note the rapid emergence of two continuous phases; that is indicative of rapid flow.

The third mechanism accounts for an intermediate droplet concentration, somewhere between 21 and 35% of the volume occupied by the droplets.  This is the most difficult mechanism and has been the subject of various researchers for decades.  That is because neither diffusion nor hydrodyanmcs can be neglected and thus makes it difficult to make approximations.  Hydrodynamic forces stem from thermally-induced Brownian motion, the technical term for the random collisions between particles in a fluid, as seen in Figure 3.    Diffusion occurs during the collision, and thus the hydrodynamics and diffusion are coupled.

Figure 3.  Animation of Brownian motion without coalescence.   The blue path traces the trajectory taken by the yellow particle from  incessant collisions.

However, what Shimizu and Tanaka claim is that the collisions are not random, but are driven by a so-called composition Marangoni force, which is induced by long-range interfacial tension gradients on the droplets.  In other words, the interfacial energy on a droplet varies because of other droplets near it.   This is illustrated from the paper in Figure 4.

Figure 4.  Snapshot of droplets with interfacial energy gradients, with blue color being the minimum, and red as the maximum.  Velocity arrows are drawn on and in the droplets of b to indicate preferred velocities on a given droplet region.  Thus, the droplet in the low left is drawn toward the small droplet next to it due to the largest arrows pointing toward it from the further end of this droplet.

A droplet much smaller than its neighbors draws neighboring droplets toward it and eventually induce them to collide.  In short, a larger droplet tends to eat a smaller droplet by direct collision.   The attraction from droplets stems from the different gradients of surface tension between the droplets.  The merging then lowers the overall energy of the system, since equilibrium just means the system has fallen to its lowest energy state.  

Source:  R. Shimizu, H. Tanaka, A novel coarsening mechanism of droplets in immiscible fluid mixtures, Nature Communications, 6 (2015).

Probing C and N doped titanium dioxide with hard x-ray photoelectron spectroscopy

Titanium dioxide (TiO₂) is a widely used band-gap semiconductor with vast applications in photocatalysis, photovoltaic, and spintronics.    Recent advances by Japanese researchers has sparked interest in TiO₂ for electro-photo-catalytic splitting of water, since this mineral is a widely used white pigment with UV absorbance characteristics.  The Achilles heal of this application is its large band-gap (3.20 and 3.00 eV for anatase and rutile crystal phases, respectively).  Hence the push toward doping for lower band-gap.  Among the several anionic dopants attempted, N was the most effective for lower band-gap and thus enhance photocatalysis under visible light.  This is made possible by the overlap of the N 2p orbitals with the O 2p orbitals with respect to the energy it would take for a photon to knock an electron out of those orbitals and into the so-called conduction band, where current can be generated.  This makes intuitive sense since N and O are neighbors on the periodic table.  Doping with N can be further enhanced with C.  Doping with only carbon is bad because its 2p orbitals do not overlap with the O 2p orbitals, which would create separate quantum levels for conducting electrons to occupy and not do anything.  However, the C 2p orbital overlaps with the N 2p orbital.  This means we'd have 2p orbitals from two dopant atoms overlap with the oxygen 2p, which gives engineers more flexibility in lowering the TiO₂ band-gap.

Up to then, little data has been collected on these C- and N-doped TiO₂. The paper by Ruzybayev et al attempts to fill those gaps.   Experimental data was acquired via hard x-rays from the National Synchrotron Light Source (NSLS)  in Brookhaven National Laboratory (now closed because NSLS II is opening now).  Hard x-rays have an energy range of 5-15 keV (about the same as medical x-rays), and penetrate into the sample bulk; soft x-rays range 1-5 keV and only penetrate the surface of a sample .  Both types of x-rays are useful, but it depends on the application.  This makes for an interesting paper because my master's thesis was a theoretical study of hard x-rays inducing photoemission from a magnetic multilayer structure.  The technique is known as hard x-ray photoelectron spectroscopy (HXPS).

The experimental section will be skipped for brevity, but the specimens were TiO₂ films that were only 500 nm thick.  That's close to the wavelength of blue light.  The optical band-gap of pure TiO₂ and co-doped TiO₂ is shown in UV-Vis diffuse reflectance spectroscopy data on Figure 1.

Figure 1.  Approximated band-gaps for pure TiO2 and C and N doped TiO2.

Pure TiO₂ has a band-gap of 3.30 eV, as evidenced by the upward slope of the solid curve in Figure 1.  Codoping lowered the band-gap to 2.39 eV, (deeply slanted slope of the dashed curve).  The analysis from HXPS data is displayed in Figure 2.  This can be considered the "meat" of the paper because it explicitly shows how the C and N dopants affect the electronic structure of TiO₂.  "Electronic structure" has many contexts, but the important one here is the oxidation state of titanium.

Figure 2.  HXPS of titanium 2p orbital in pure form and doped with C and N.

The two peaks in the pure TiO₂ spectrum show the characteristic 2p_3/2 and 2p_1/2 spin-orbit split of Ti⁴⁺ (marked A and B in the upper curve, respectively).  The superscript 4+ represents its oxidation state, i.e., the charge experienced by the titanium atom after giving away four electrons to two oxygen atoms (each oxygen atom having two additional electrons).  Doping induced two additional peaks adjacent to the the orbitals.  The 'C' and 'D' peaks represent Ti³⁺ due to the extra net electron contributed by the dopants, which then produces oxygen vacancies.  The vacancies produce an occupied Ti 3d orbital just above the valance band maximum, which is the last quantum state for an electron to occupy before surpassing the band-gap to reach the conduction band.  The jump to the conduction band is the key to generating electricity from light.  The smaller the jump for the electron, the easier it is to generate electricity from a hypothetical TiO₂ cell.

Photoelectron spectroscopy can be measured at ultra-violet wavelengths.  That is useful for measuring the valance band maximum, which the authors have done here.  More specifically, they measured the change in photoelectron kinetic energy relative to the O 2p orbital.

Figure 3.  Valence orbital data of pure, C doped, N doped, and both C and N co-doped TiO₂.  These spectra are measured relative to the O 2p orbital.

Spectra from TiO₂ doped only with N or only with C are included to ascertain the contributions to the altered band-gap by the individual dopants. The key feature is the tailing of the large peak at 5 eV.  The curve for pure TiO₂ is flat at this kinetic energy, but doped TiO₂ curves are still still slanting downward here, with the C and N co-doped curve being the highest one; in jargon, the C and N co-doped curve has the highest valence band maximum.

Experimental data was compared with data calculated from computational models that varied the locations of dopant atoms in a TiO₂ unit cell.  According to Figure 4, what matters is which atoms are the dopant atoms bonded to.  That in turn affects the photoelectron spectra due to the so-called density of states (DOS), which I won't go into because that is too advanced for this blog.  What I can say is the DOS allows you to calculate photoelectron spectra that you then compare to experiment.  If theory  and experiment don't match, try a different model.

Figure 4.  Density of states for co-doped TiO₂ unit cells.  In the insets, blue, light blue, red, and orange spheres represent Ti, O, C, and N atoms, respectively.

Trial-and-error led the authors to conclude that the model in the lower left of Figure 4 was the closest match to the spectra from the experiment, as seen in Figure 5.

Figure 5.  DOS and experimental photoelectron valence band for C and N co-doped TiO₂.  The experimental curve is the green curve from Figure 3.  The red curves are the theoretical photoelectron spectra, and the dark curves (excluding experiment) are the DOS.

That unit cell is represented in the middle spectrum of Figure 5.  It is a fairly good match despite the sloping background of the experimental spectrum.  This has led the authors to conclude that carbon preferentially sticks to titanium, while nitrogen prefers oxygen.  The results show that electronic structure of TiO₂ can be manipulated to decrease the band-gap for photocatalysis.

Creation of triblock copolymer thin films by combining vapor annealing with a raster spray

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Adapted from "Spatial and Orientation Control of Cylindrical Nanostructures in ABA Triblock Copolymer Thin Films by Raster Solvent Vapor Annealing", ACSNano,

Jonathan E. Seppala, Ronald L. Lewis, III, and Thomas H. Epps, VOL. 6 ’ NO. 11 ’ 9855–9862 ’ 2012)

This nanosciece paper describes a novel approach to annealing polymeric thin films, particularly block copolymer thin films.  Self-assembly is a significant phenomenon in these materials because they open the floodgates for designer nanoscale materials for nanoporous membranes, lithographic masks, and nanopatterning/templating applications (the last two have huge implications for the electronics industry).  These three nanotechnologies often exploit morphologies often found in AB diblock and ABA triblock copolymers (spheres, gyroid, and lamellae) because the thermodynamics of bulk self-assembly is relatively well established.    The morphology of bulk block copolymers is influenced by three major factors: the degree of polymerization (N), the interaction parameter (c), and the volume fraction of the blocks (f). 

With thin films, surface energy becomes an additional factor; it can be exploited by thermal annealing to facilitate copolymer self-assembly via the bestowment of mobility to amorphous regions that are trapped upon casting.  However, thermal annealing is limited to copolymer systems where the components have similar γ‘s and are thermally insensitive.  Another technique is solvent vapor annealing (SVA), which grants mobility by effectively reducing the T_g of the copolymers.  It is a powerful technique, but is limited to small quantities that are typical of research labs. 

The objective is to devise a faster method for large-scale production of block copolymer (BCP) nanotechnologies.  The method must enable control over morphology and orientation of BCP thin films.  The authors propose raster solvent vapor annealing (RSVA): solvent vapor from a bubbler system is directed onto a BCP surface, which then creates a SVA zone.  The zone is modified/expanded by a motorized stage moving in a raster fashion.  

RSVA was performed with a THF-rich vapor stream in single or multiple passes over a 100 nm thick poly(styrene-b-isoprene-b-styrene) (SIS) film, with domains of 29 nm.  The RSVA speeds ranged from 500 µm/s to 3 µm/s.   The RSVA process swelled the films due to hydrolysis, so the film thickness was measured by spectral reflectometry.  The swelling increased the thickness to 160 nm, and eventually dried down to nearly the original thickness, although some samples were reported to have residual solvent. 

Several approaches to RSVA were performed.  One was single-pass, where the stage moved under the nozzle once, with only the speed varying. The as-cast film had a lamellar structure consisting of cylinders oriented parallel to the substrate, but with minimal long-range order.   Varying the speed affected the ordering of the lamellar cylinders.  The slower the speed, the longer the order-range; at 10 µm/s, the cylinders have mostly perpendicular orientation.  The 10 µm/s speed corresponds to an annealing time of 50 s.  The cylinders even looked slightly swollen, which was confirmed by azimuthally integrated 1D profiles from FFTs of the AFM images.    Even reducing the nozzle diameter still induced the ^cylinders, although at lower speeds. 

           

The post-RSVA morphology is an imbroglio of competing forces. The lower g for polyisoprene (32.0 mJ/m² vs. 40.7 mJ/m²), the majority block, enables wetting of both the free and substrate surfaces, which leads to the propensity for the cylinders to possess parallel orientation.  High RSVA speeds do not change the orientation (see 2a–2c) because the surface energy difference was too large for entropy to take effect.  The slow raster allowed enough solvation to lower the differences in g, which lets a) entropic effects to manifest, and to compensate for the stretching experienced by the cylinders during swelling and deswelling. 

              Briefly, Seppela et al tried two more approaches.  One is multiple passes under the nozzle.  Retracing the RSVA pathways altered the cylindrical orientation toward perpendicularity.  The other is a crossed-path approach; two orthogonal passes cross each other, and the result is a domain dominated by perpendicular cylinders.  Both approaches are supported by crisp AFM phase images. 

            This is a wonderfully written paper, but it helps that I have significant background in polymer chemistry.  Note that only THF was used in the vapor stream; it’s natural to ask if this approach has been done with other solvents, and other BCPs.  d_THF = 18.1 (MPa)^1/2, d_polyisoprene = 16.2 (MPa)^1/2, d_polystyrene =  18.6 (MPa)^1/2, so utilizing similar solubilities is a probable reason for the RSVA setup described in the paper.  One should also ask if this technique can be done for other self-assembling polymeric thin films. Seppela et al noted that this annealing method can be altered according to slit geometries, solvent quality, and substrate temperature–indications of much promise for RSVA.  

Confined Crystallization of Polyethylene Oxide in Nanolayer Assemblies

We live in an era of increasing reliance on the very small to satisfy humanity’s endless needs and desires for new technologies.  Nanotechnology manifests itself in numerous scientific fields, and polymer chemistry is no exception.  Polymers are generally amorphous, but polymer crystallinity can be observed if the conditions are right.  Semi-crystalline polymer chains (possesses crystalline and amorphous phases) such as polyethylene and nylon are often used as barrier films in food, medicine, and electronics industries.  A barrier is considered highly efficient if small gas molecules are relegated to permeating through only the amorphous regions of the chains (crystalline regions are impenetrable).   Efficiency can be fine-tuned by varying the polymer-film processing conditions to suit the desired amount of crystallinity and chain orientation.  Polymer films can now be made thin enough to effectively confine the crystallization process to 2D; this leads to surprising results. 

Conventionally, confined polymer chains crystallize into lamellae with thicknesses of ~10-20 nm with spherelitic morphology.   However, this convention is skirted at the nanoscale, as isotropic growth is severely hampered to the point of producing lamellar crystal orientation.  This orientation is usually perpendicular to the layer (edge-on), but parallel orientations have been reported several times in the literature; mechanisms for orientation determination remain mysterious for the time being. 

Normally, researchers prepare 2D crystallization of polymers via solution processes such as spin-coating or Langmuir-Blodgett (LB) techniques, but these are limited by the solvent requirement and the small quantity of material fabricated.  LB techniques enable layered nm morphologies due to microphase separation of dissimilar block copolymers within the thin films.  Alas, block copolymers are notoriously difficult to synthesize and align with respect to the direction of the thin films. 

Enter a new technique known as layer-multiplying extrusion.  It uses forced assembly to create alternating layers of two polymers that number up to the 100,000s.  Almost any melt-processable polymer can be formulated into kilometers of nanolayered films with thicknesses of ~10 nm.  With less material comes an explosion of new previously unknown properties (“less is more”). 

The materials used in this study are polyethylene oxide (PEO, also known as polyethylene glycol), which has the following structure:

                                                   HO-CH₂-(CH₂-O-CH₂-)_n-CH₂-OH

The other is ethylene-co-acrylic acid (EAA), a copolymer with much lower crystallinity than PEO:  

Films with 33, 257, and 1025 alternating EAA and PEO layers were extruded, with various thicknesses and composition ratios, including (EAA/PEO vol/vol) 50/50, 70/30, 80/20, and 90/10.  The nominal PEO layer varied from 3.6 µm to 8 nm. 

The films were subjected to oxygen permeability tests with respect to to layer thickness.  The results are shown below:

Fig. 1 The effect of layer thickness on oxygen permeability. (A) Oxygen permeability of films with equal volume fractions of EAA and PEO. The dashed line indicates P_// calculated from Eq. 1. (B) Oxygen permeability of the PEO layers from films of varying composition calculated from Eq. 2. The dashed line indicates PPEO. The open symbol is for a film with PEO layer breakup. The solid lines are drawn to guide the eyes. 

The plots show a significant decrease in O₂ permeability.  Gas permeability for layered assemblies is modeled by the following equation. 

     (1)

where 𝜙_PEO is the volume fraction of PEO and P_PEO and P_EAA are the permeabilities of PEO and EAA, respectively.  Upon plugging determined values of P_PEO and P_EAA from literature into Eq. (1), the result did not agree with the findings reported in the plot above. Eq. (1) predicts increasing permeability with respect to decreasing PEO thickness, but the data show the opposite trend. Eq. (1) was then modified to account for the apparent sensitivity to P_PEO due to the far lesser permeability of PEO; it still did not agree with the plotted data with the exception of thicker PEO layers as indicated by the dashed line.  Clearly, the PEO nanolayers possess some previously unknown crystalline morphology that bestowed them with staggeringly low permeability.  However, differential scanning calorimetry revealed that the PEO and EAA layers (even the very thin ones) share the same melting enthalpy and melting temperature as the control films; this means that the changes in crystalline morphology granting the PEO nanolayers low permeability was not accompanied by changes in crystallinity nor lamellar thickness. 

Upon examination by AFM, the authors found that the thin 20 nm PEO layers exhibited single lamellae that extended beyond the field of the AFM image.  The single lamellae are said to be very large single crystals.  Reducing the PEO layer thickness to 8 nm then induces breakage, thereby increasing the permeability.  Fig. 2 below shows the AFM image of the 20 nm PEO layer, and an accompanying schematic showing a gas diffusion pathway through the layered assembly.

Fig. 2  AFM phase images of partial cross sections of the layered EAA/PEO films. The PEO layer has substantially higher crystallinity than the EAA layers and hence appears bright in the AFM images. (A) A low-resolution image of an EEA/PEO film with 50/50 composition, 33 alternating layers, and nominal PEO layer thickness of 3.6 mm. (B) A higher-resolution image showing the spherulitic morphology of the 3.6-mm-thick PEO layer. (C) A low-resolution image of an EAA/PEO film with 70/30 composition, 1025 alternating layers and nominal PEO layer thickness of 110 nm. (D) A higher-resolution image of the 110-nm-thick PEO layers showing the oriented stacks of PEO lamellae. (E) A high-resolution image of an EAA/PEO film with 90/10 composition, 1025 alternating layers, and nominal PEO layer thickness of 20 nm showing that the PEO layers crystallized as single, extremely large lamellae. (F) A schematic showing the gas diffusion pathway through the layered assembly with 20-nm- thick PEO layers. The arrows identify the EAA layers and PEO layers. 

The lamellar crystalline region is considered impermeable, with the lamellar fold surfaces constituting the permeable amorphous regions.  As seen in Fig. 2, the gas pathways depend on the frequency of defects such as lamellar edges.  The permeability is now expressed by

   (2)

 where α is the aspect ratio of the impermeable platelets (length/width), and 𝜙 is the volume fraction of impermeable platelets; the platelets are orientated perpendicular to the flux.  For the thinnest PEO layers, the aspect ratio was as high as 120, which meant the lamellae extended up to 2 µm for the 20 nm thick layers.   Gradually thickening the PEO layer relaxed the restrictions on 3D growth, which returned the morphology to spherelitic.   The results were further confirmed by small-angle x-ray scattering (SAXS) and wide-angle x-ray scattering (WAXS). 

This work is a major breakthrough in polymeric applications for nanotechnology because it shows experiment trumping theory, and possibly describes a major advance for gas-barrier films.  Its importance is amply demonstrated by the 51 citations it has generated since its publication in 2009.  Science Magazine accepted the paper because of its reliance on well-established analytical techniques (AFM, differential scanning calorimetry, SAXS, WAXS), and, more importantly, because of its broad significance in the field of nanoscience. 

This significance is underscored by the novel utilization of a relatively new technique–coextrusion–on readily available polymers to engineer nanolayered polymeric formations in sufficient amounts to allow for probing links between the confined crystalline morphology and the properties exhibited.  This opens up new possibilities for packaging methods, i.e., incorporating polymer nanolayers into common polymeric films for less cost, thereby reducing the environmental and energy consequences.  

Laser Ablation: Discussion & Conclusion

Fig. 11 from the last entry does not bode well for laser ablation as a profiling technique for CARCs (chemical agent resistant coatings).  Why?  Because it didn't resolve the UV-damanged region in the topcoat.  At least this was a feasibility study, so its purpose was fulfilled, but a better alternative to ATR-mode FTIR depth profiling still awaits discovery.  Fig. 6 from the last entry shows remarkable resilience from the signature peaks after ablation.  That should mean then ablation shouldn't be a factor when one investigates the coating after QUV exposure (accelerated weatherization under controlled conditions).  Fig. 1 shows why.

Fig. 1 FTIR spectra of major organic and inorganic bands for baseline sample (1), 15-mm-deep transmission-mode spectrum in UV-aged sample (2) and 15-mm-deep ablation window in UV-aged sample (3). 

There shouldn't be discernible differences for spectra (2) and (3), but difference is obvious for the carbonyl peak on the left.  Considering that most CARCs (to my limited knowledge) have polyurethane binders, this can be considered a death blow to the possibility of laser ablation being used as a depth-profiling technique for CARC (chemical agent resistant coating) films after long-term exposure to the elements.  The authors speculate that the ablation process creates ether groups (C–O–C), which overlap with carbonyl groups.

In addition, the amide peaks seen in the spectra of the aged samples likely stem from other functional groups that overlap; they might result from a complex interaction between the aging and ablation processes.  There's still the chance that the original amide II group had reformed after the ablation, which explains the awful ablation profile in Fig. 10(b) of the last blog entry.  This reformation effect was seen in previous studies involving UV-induced cross linking between proteins and DNA with little disruption to the bulk protein chemistry.  This is important considering the chemical similarities between peptide bonds (–CO–NH–) in protein and urethane bonds (–O–CO–NH–) within polyurethane.  

The greater activity within the carbonyl region of Fig. 1 above is perhaps caused by a carboxyl group  (–CO–OH) rather than carbonyl or even ether.  If so, there should be larger peaks in the –OH stretching region (~3300 cm^-1), but not so large that it surprises the aged-but-unblated sample.  Alas, Fig. 2 below shows this is not the case.

Fig. 2  FTIR spectra showing (OeH) and (CH2) bands for baseline sample (1), 15-mm-deep transmission-mode spectrum in UV-aged sample (2) and 15-mm-deep ablation window in UV-aged sample (3). 

(3) lies between (2) in the CH2 stretching region (2937 cm^-1), but not in the -OH area (3364 cm^-1).  They attribute this to another unforeseen reaction with the ablation process.  Nevertheless, it's clear from here that femtosecond laser ablation is unreliable as a depth-profile technique for aged CARC films.

Laser Ablation: Results

The figure below shows the ablation window depth as a function of stage scan speed.  The removal rate shows exponential decay, but the obvious fact points towed lower removal rates at lower energies; what's important is that the user can still control how stuff is removed by laser.

 Fig.1 Calibration profiles for three different energy levels at several scanning speeds (dwell times). Note absence of ablation window at 35 mJ energy level regardless of scan speed. 

Fig. 1 shown below, shows signs of a minimum window depth feasible within any pulse energy given the maximum scanning speed available. The most important curve is the seemingly flat 35 µJ pulse energy profile, (1 µJ = 10^-6J) which stays flat no matter the scan speed, as seen above.  The other two energies (167 and 73 µJ) show clear decay slopes, which can be interpreted as progressive erosion by sample speed reduction (going backwards on the x-axis in Fig. 1 above).  The calibration profiles on the left conform to this finding.

 Fig.2 SEM micrograph showing surface detail of baseline (unablated) coating sample.    

Figs 2-4 are scanning electron microscope (SEM) micrographs, AKA, closeups, of the coating after different laser energies.  Fig. 2 is unabated; Fig. 3 depicts the coating after the 35 µJ laser; Fig. 4 is after 167 µJ.  The coating itself is a TiO₂ pigment (small white particles) embedded within the polyurethane binder (uniform dark background); the large spheres and blocky objects are siliceous and talc fillers, respectively.  

Fig. 3 SEM micrograph showing surface detail of coating sample after 35 mJ ablation, 10K mm/s scan speed. Note selective ablation of TiO2 pigment (evidenced by presence of ovaloid cavities) with no apparent disruption to surrounding organic binder. 

Fig.4 SEM micrograph showing surface detail of coating sample after 167 mJ ablation, 10K mm/s scan speed.    

 The laser darkens the pigment to various degrees, even for the 35 µJ laser, which did not ablate the coating.  The 167 µJ scan destroyed most of the pigment and considerably altered the binder morphology to the point of creating a large crater on the surface.  Fig. 5 is a cross-section of the coating and shows a sharp distinction between the ablated and non-ablated regions.

Fig. 5 SEM micrograph of paint cross-section showing transition between unablated and ablated surfaces. Boxed areas are shown in detail in the lower two images. On the left is detail from the unablated surface showing uniform distribution of pigment (small white particles). The right image shows detail from the ablated surface. A 5–8 µm pigment-depleted layer is indicated by the red line. 

It's clear from Fig. 5 that the ablated region has lost much of the pigment and differs significantly in morphology.  TiO₂ is a semiconductor with a relatively small band gap which makes it more susceptible to ablation than the surrounding polyurethane binder.

It's going to get a little technical from here on out, but this is where it gets quantitative.  Fig. 6 shows the transmission spectra of a baseline (pristine) sample plus spectra from all three ablation energies (35, 73, 173 µJ).

Fig. 6. Transmission-mode spectrum of unaged baseline and ATR-mode spectra for 35, 73 and 167 µJ ablation. The ester and urethane peaks experienced little change from the ablation process.

The peaks contained within the box on the left did not change much after ablation; this is to be expected since polyurethane binder is mostly what's left in the ablated regions.  There's a large gain in the peak on the right (indicates a gain in quantity for that functional group), but that's a region of considerable overlap between the C–O–C ether and O–Si–O; the authors guess it to be Si–O, since there's a lot of new SiO₂ filler exposed.  


Now I compare samples that differ by the amount of simulated weather that they've been exposed to.  Below is Fig. 7.

Fig. 7 ATR-FTIR spectra for baseline, 6-week and 18-week QUV exposed samples of coating. Near-extinction of amide II peak (1523 cm^-1) is noted, evidence of photooxidation of urethane groups in the organic binder of the coating. (b) OH absorption (3364 cm^-1) is seen to broaden substantially, whereas hydrocarbon peaks (2937, 2863 cm^-1) simultaneously decrease, as a result of QUV exposure. 

Keep in mind that cm^-1 (wavenumbers) is a unit of frequency just like Hz (s^-1); in fact it's proportional by E = ℎν = ℎc/λ for the infrared radiation (IR).  You could see from Fig. 7 that the IR is compared for the baseline, 6-weeks QUV exposed, and 18-weeks QUV exposed.  The authors emphasize the complete extinction of the so-called amide II functional group, which means drastic decay for the polyurethane binder.  In part (b) of Fig. 7, the OH/NH band broadens, perhaps due to carboxyl group formation.  Simultaneously,the QUV process reduces the hydrocarbon peaks in the region close to 3000 cm^-1.

Fig. 8 presents depth profiles of the ablated window for the baseline and 18-week QUV samples.

 Fig. 9 (a) Ablation window depths for 140 µJ pulse energy, baseline and 18-week QUV samples. (b) Ablation window depths for 80 µJ pulse energy, baseline and 18-week QUV samples. At all scan speed/ pulse energy combinations material removal rate for UV-aged sample exceeds baseline sample 

The depth windows show a greater rate of material removal for the QUV-aged sample than for the baseline; probably due to the weathered urethane binder becoming more broken down.

Fig. 10 shows FTIR depth-profile measurements.  This is the nitty-gritty of the study, as it shows why this method falls short.  

Fig. 11 a) Amide II/C]O ratio for baseline vs. UV-aged coating samples using cross-section transmission-mode FTIR. UV-zone extending 30 mm into topcoat is readily apparent. (b) Amide II/ C]O ratio for baseline vs. UV-aged coating samples using femtosec- ond ablation-assisted depth profiling. UV-damaged zone not apparent. 

Fig. 11 a) repeats a result from an older study where they used transmission-mode FTIR on 3 µm thick cross-sections of the same two samples.  It's mentioned here for the sake of comparison.  Peak ratios are used here to eliminate fluctuating values due to varying sample thicknesses.  Part (a) shows a clear change in the ratios as the depth is increased, meaning the near-surface region is more weather-damaged than the deeper regions in the bulk.  That is not the case for part (b); if the ablation method did its job, it would show a damage gradient just like in part (a).  The next blog will discuss why this is so.