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Multiphase coatings from complex radiation curable polyurethane dispersions
Progress in Organic Coatings 75 (2012) 560–568
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Progress in Organic Coatings journal homepage: www.elsevier.com/locate/porgcoat
Multiphase coatings from complex radiation curable polyurethane dispersions Michel Tielemans a,∗ , Patrice Roose a , Chinh Ngo b , Roberto Lazzaroni b , Philippe Leclère b,∗∗ a b
R&D, Cytec Surface Specialties, Anderlechtstraat 33, B-1620 Drogenbos, Belgium Laboratory for Chemistry of Novel Materials, University of Mons, Place du Parc 20, B-7000 Mons, Belgium
a r t i c l e
i n f o
Article history: Received 10 August 2011 Received in revised form 25 April 2012 Accepted 22 May 2012 Available online 15 June 2012 Keywords: Polyurethane Waterborne Radiation-curing Morphology Film-formation Atomic force microscopy Pigments
a b s t r a c t The waterborne nature of radiation curable polyurethane dispersions largely respond to the current environmental concerns and do not require any additional coalescent since the film formation (drying) and hardening (photo-curing) take place in distinct steps. It is possible to design aqueous dispersions with distinct polymer particle populations resulting in micro-structured coatings with optimized properties over a wide range of curing conditions. Mixed dispersions based on hard and soft acrylated polyurethane particles were used as model systems for the present study. The minimum film formation temperature has been investigated as a function of the hard:soft polymer ratio. The elastic modulus of the dry coatings shows a reinforcing effect consistent with the inclusion of hard domains in a soft continuous matrix. However, the level of reinforcement is not properly predicted by the usual mechanical models and it is qualitatively accounted for by assuming a composition gradient (interphase) between the hard domains and the matrix. The multiple-phase structure was clearly established by Atomic Force Microscopy in agreement with thermal analysis data. Furthermore the local nanoscale mechanical properties were mapped using a new imaging mode based on real-time force–distance curve analysis. Finally, the coatings prepared using this multiple-phase pattern present a clear benefit over conventional homogeneous coatings by offering an improved balance of chemical and mechanical resistance in pigmented systems applied on melamine-coated MDF panels. © 2012 Elsevier B.V. All rights reserved.
1. Introduction The introduction of water-based alternatives profoundly changed the landscape of the radiation curing market by combining the high productivity and performance of traditional radiation curing compositions with the low viscosity of usual water-based systems [1]. Aqueous colloidal dispersions of acrylated polyurethane oligomers (referred to as UV-PUD’s) largely respond to the environmental regulations due to their waterborne nature and the absence of volatile organic compounds or additional coalescents, since the film formation (drying) and hardening (photo-curing) take place in distinct steps. They have a good colloidal stability [2] and their rheology [3] makes them particularly suitable for application by spray, curtain or roller onto different substrates. A wide range of products has been developed for clearcoat application on wood furniture with excellent adhesion, chemical and mechanical resistance for indoor [4] as well as outdoor [5] applications.
∗ Corresponding author. ∗∗ Corresponding author. E-mail addresses: [email protected] (M. Tielemans), [email protected] (P. Leclère). 0300-9440/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.porgcoat.2012.05.010
The coatings obtained after film formation and photocuring are characterized by exceptional mechanical properties since they combine the usual features of polyurethanes along with high cross-linking levels. The microstructure of the cured coating is typically characterized by an alternation of hard and soft segments entrapped within a dense crosslinked network, providing a particular ability to balance contrasting properties, such as flexibility and adhesion, to various substrates on one side, scratch resistance and chemical resistance on the other side. As long as clearcoats are concerned, light penetration and in-depth curing (“through cure”) of the coating is not an issue. However, a technological limitation arises when the radiation curable polyurethane dispersions are formulated with organic or inorganic pigments so as to obtain a colored decorative coating. In this case, substantial ultraviolet light is absorbed and scattered by the pigments which prevents efficient light penetration into the coating, in particular for pigments other than white. In order to design a colored coating with sufficient mechanical strength in the bulk along with strong chemical resistance at the surface, without sacrificing the film formation and the adhesion at the coating–substrate interface, a mixing approach with tuned particle characteristics has been suggested recently [6]. It is possible to make from soft to very hard polymer particles dispersed in water and presenting each the characteristics inherent to their chemical nature. Soft particles will provide low
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other existing and promising AFM techniques such as AFM Noise Analysis [12] or Band Excitation AFM [13]. Finally, the product has been formulated with a high pigment load and applied by spray on melamine-coated MDF panels prior to the curing under UV light. The coatings prepared using a multiplephase structure of hard and soft acrylated polymeric components present a distinctive benefit over conventional homogeneous coatings. 2. Materials and methods
Fig. 1. Schematic representation of polymer dispersions made from (a) soft particles, (b) hard particles, (c) hard core–soft shell particles and (d) soft and hard particles.
film formation temperatures without any coalescing solvent but the mechanical properties of the bulk film before cure (or in the absence of cure in case of strong light interaction with pigments) are very poor. In contrast, hard particles will enhance the mechanical properties before cure but the high film formation temperature, requesting significant levels of coalescent, is totally inappropriate and usually associated with a poor adhesion and depressed optical properties. It appears that only complex polymer dispersion morphologies are suited to balance the coating properties in the appropriate manner. Although core-shell morphologies offer some real attractiveness for designing multiple-phase coatings, they are much more difficult to control and to manufacture so that their benefits over particle blending are not always straightforward. The dispersions originating from soft particles, hard particles, hard core–soft shell particles and a physical blend of soft and hard particles are schematically depicted in Fig. 1. The UV-PUD system used in the present study is based on a blend of hard (i.e., with a glass transition temperature, Tg , higher than room temperature) and soft (with Tg lower than room temperature) acrylated aliphatic polyurethane particles and demonstrates that well balanced properties can be achieved before and after UV-curing following this approach. In this paper, we first examine the minimum film formation temperature (MFFT) as a function of the hard:soft ratio in the blend composition. Next, in order to characterize the phase structure of the coatings, the thermal and thermo-mechanical properties are measured by temperature-modulated differential scanning calorimetry (TMDSC) and thermo-mechanical analysis (TMA). The tensile properties show clearly a reinforcing effect of hard domains embedded in a continuous matrix. Many empirical or semi-empirical models have been proposed to predict the modulus of particulate–polymer composites [7–11]. However, the assumption of a composition gradient (“interphase”) between the hard inclusions and the soft matrix was required in order to capture the enhancement of Young’s modulus. The results have been correlated with the microphase characteristics determined by Atomic Force Microscopy (AFM). Moreover, the local mechanical properties, such as stiffness, adhesion, dissipation and deformation, were mapped by using a new imaging mode based on real-time force–distance curve analysis. This recent technique, referred to as Peak Force Quantitative Nanomechanical Property Mapping (PFQNM), returns local mechanical properties at the nanoscale within reasonable acquisition times compared to
UV-PUD’s are anionically stabilized acrylated polyurethane colloids in water [14]. A typical polyurethane ionomer structure is composed from (i) a polyacrylated molecule having at least one chemical functionality able to react with isocyanates; (ii) a polyisocyanate; (iii) a polyol selected among polyesters, polyethers or polycarbonates; (iv) a functional polyol capable to disperse the oligomer in water and whose functionality is usually a carboxylic acid or a salt thereof and (v) a polyamine (optional). They encompass a very large range of composition and molar mass with a linear or branched polymeric architecture and a varying amount of urethane, urea, allophanate and biuret functionality. The preparation of the urethane–acrylate polymer follows a multi-step process in a low boiling point solvent (e.g. acetone). The resulting polymer solution is then neutralized, dispersed and optionally chain-extended in water, the solvent being finally removed from the dispersion medium under reduced pressure. Two base dispersions were prepared for the purpose of this study, each originating from the same carboxylated and acrylated polyurethane backbone but clearly differentiated by their molar mass distributions. The first model dispersion is a soft (low Tg ) polymer dispersion with a low molar mass and a minimum film formation temperature (MFFT) below ambient temperature. The second model dispersion is a hard polymer dispersion with a very high molar mass range and a MFFT much higher than room temperature. The main characteristics of the two model dispersions are summarized in Table 1. The solid content of the dispersions was determined gravimetrically after 2 h at 120 ◦ C. The viscosity was measured with a Brookfield viscometer at 25 ◦ C and at a rotational speed of 50 rpm. The hydrodynamic particle diameter (dDLS ) was determined by dynamic light scattering using a Malvern Autosizer LoC instrument linked to a Malvern Series 7032 Multi-8 correlator. The colloidal stability was assessed by multiple light scattering (TurbiscanTM ) at 60 ◦ C by recording the number of days before critical destabilization. The weight-average molar mass (Mw ) was measured by gel permeation chromatography on the soluble polymer fraction in tetrahydrofuran. The two base dispersions have been investigated separately or mixed as blends in various ratios in order to characterize the film formation properties and to relate the microstructure to the mechanical properties in the dry state after film formation. A series of blends were prepared according to the hard:soft polymer ratio of 20:80, 30:70 and 40:60. When required, a photo-reactive
Table 1 Characteristics of the proprietary polyurethane dispersions used for the blends.
Solids (%, w/w) Viscosity (mPa·s) pH dDLS (nm) Stability at 60 ◦ C (days) Mw (kDa) Solubility in THF (%)
Soft
Hard
35 100 7.5 70 >10 15 100
35 100 8.0 60 >10 >100 ∼15
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formulation was obtained by adding 1.5% (w/w) ESACURETM HB as photoinitiator to the dispersions. The MFFT was measured according to ASTM 2354 on a Rhopoint® automatic gradient-heated metal plate. The film formation kinetics of the dispersions has been addressed using adaptive speckle imaging interferometry (ASII) by applying coatings (100 m) on glass at 22 ◦ C [15]. This technique is based on laser light scattering of a drying film where the resulting interference image is recorded with a camera detector. The interference pattern (speckle) changes with the reduction in mobility of the scattering elements and eventually provides a qualitative characterization of the film formation over time. Atomic Force Microscopy (AFM) is increasingly used for performing mechanical studies of materials on the nanometer scale. Widely known are single molecule force spectroscopy [16], shear force modulation [17], friction microscopy [18], force modulation [19] and slightly less quantitative, but providing pertinent information, the intermittent contact or Tapping Mode [20]. For the morphological characterization of the film microstructure, Tapping-Mode AFM (TM-AFM) was used on a Veeco (now Bruker Nano Inc.) ICON AFM (Santa Barbara, CA) driven by a Nanoscope V control unit. We operated in air and in ambient conditions (temperature and pressure). Thin coating films were prepared on cleaned glass substrates using a bar coater and subsequently dried at 40 or 80 ◦ C. The probe has a resonance frequency of about 300 kHz. At this point it is worth mentioning that, in order to keep a well-defined oscillating behavior of the tip, the perturbation to the oscillator due to the contact with the surface is chosen to be small. In other words, the reduction with respect to the free amplitude (the setpoint) is only of a few percent (typically 5%). This method has two advantages. From an experimental point of view, this allows to identify easily hard and soft areas, the strongly phase-shifted areas of the image corresponding to the hard domains. From a theoretical point of view, it permits to use simple approximations providing analytical solutions that help to interpret the experimental data. Free-standing coating films were prepared for the determination of the thermo-mechanical properties of the different systems. The aqueous UV-PUD resins were applied on an untreated bioriented polypropylene (BOPP) plastic film (30 m) using a bar coater. In order to obtain films of suitable thickness, five layers of ≈35 m (wet) were applied on top of each other and every layer was dried for 1 h in an air-convection oven. The effect of the drying temperature was studied at 40 and 80 ◦ C. The liquid polymer dispersion showed good wetting on the BOPP film and the coatings could easily be removed from the substrates after drying. A portion of every film was separated and cured twice on a UV-conveyor belt using a high power medium-pressure mercury lamp (power input 80 W cm−1 ; belt speed 5 m min−1 ). The phase behavior of the dried coatings was studied using temperature-modulated differential scanning calorimetry (TMDSC). The measurements were conducted using a Mettler DSC823e instrument with standard aluminum crucibles and nitrogen as the flow gas (50 mL min−1 ). The sample mass was typically ≈10 mg. The sample was heated using a sinusoidal temperature modulation with a period of 60 s and an amplitude of 1 ◦ C superimposed on a linear ramp of 2 ◦ C min−1 . Prior to the modulated heating step, the sample was heated at 85 ◦ C for 5 min in order to improve samplepan contact. The glass transition temperature is determined at the midpoint of the transition. Thermo-mechanical analysis (TMA) was conducted in static penetration mode using a flat-ended indentation rod (radius 1.27 mm) under a static normal load of 0.5 N. TMA data were recorded with a 2940 thermo-mechanical analyzer (TA instruments) at a heating rate of 3 ◦ C min−1 . Tensile properties were measured at room temperature with a Zwick Z010 elongation testing machine at a cross-head speed
of 1 mm min−1 . For the tests, 3 cm × 0.5 cm rectangular-shaped specimens were cut from the coatings. At least five independent measurements were performed for each sample.
3. Results and discussion 3.1. Film formation The dispersions prepared from the soft and hard acrylated polyurethanes were investigated for MFFT. The soft polymer dispersion has a MFFT value lower than 0 ◦ C as common for this kind of composition. On the other hand, the MFFT value of the hard polymer dispersion is beyond 90 ◦ C. The basic idea behind our approach is that the non film-forming component can be incorporated through the presence of the film-forming component. The MFFT has been investigated as a function of the hard:soft polymer ratio. The results are plotted in Fig. 2. The data present a gradual but slow increase of the MFFT up to ≈50% (w/w) of hard polymer (i.e., hard:soft ratio = 50:50) and then exhibit a marked rise after the transition point. This observation is consistent with a soft polymer matrix surrounding the hard particles up to the critical volume concentration where there is not enough of the soft polymer left to fill effectively the space between the hard polymer particles. Colombini et al. showed that the particle size ratio of hard and soft latex blends also has a large influence on the MFFT behavior [21]. They found that for a blend with a 50:50 hard:soft polymer ratio, the MFFT does not depart very much from the value of the soft matrix as long as the particle size ratio dsoft /dhard stays below 1.5. In agreement, the MFFT of the 50:50 blend in Fig. 2 has raised by 12 ◦ C compared to the pure soft polymer but still suggests that the continuum is essentially formed by the soft compound. Beyond this point, the formation of hard particle aggregates and later percolation leads to bicontinuous and inverse morphologies with a marked enhancement of the MFFT as a consequence [21]. For all the compositions further analyzed in this study, the soft polymer is thus effectively embedding the hard polymer by creating a continuous film forming matrix with hard polymer inclusions as observed by AFM (vide infra). The ASII experiments conducted for the soft polymer dispersion and the blend with a hard:soft ratio of 40:60 support this statement since no significant differences could be recorded from the shape of the drying profiles, showing similar characteristic drying times at 22 ◦ C.
Fig. 2. Minimum film formation temperature of coating blends as a function of hard polyurethane content (%, w/w).
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temperature of 40 ◦ C. The thin film of the pure soft polymer is flat with small, shallow holes; the surface is characterized by a rootmean-square roughness of only 0.35 nm. The roughness increases to 3.25 nm and 8.35 nm when incorporating the hard polymer at 20% (w/w) and 40% (w/w), respectively. The morphology observed on the height image (c), (e) and (g) corresponds to a dispersion of small spherical objects appearing harder than the surrounding matrix on the corresponding phase images (d), (f) and (h). From the image analysis, we can also conclude that the number of these spherical objects increases with the volume fraction of hard polymer present in the blend. For the 40:60 ratio (images (g) and (h)), the spheres are packed very close to each other. The apparent average diameter of the spherical objects is also slightly increasing with the hard:soft ratio (56 nm for 20:80, 58 nm for 30:70 and 60 nm for 40:60). The samples dried at 80 ◦ C (images not shown here) are very similar in terms of morphology and average diameter of the hard domains (58 nm for 20:80, 62 nm for 30:70 and 65 nm for 40:60). From these results, we can conclude that the hard polymer is well micro-phase separated from the soft polymer matrix and forms spherical objects visible in TM-AFM images with increasing apparent diameters relative to the hard:soft ratio. Despite the success of TM-AFM for micro-structural characterization, important questions remain about the physical origin of the image contrast. The height images are generally considered to display topographic information, but it must be kept in mind that the local mechanical properties of the sample (i.e., the possibility that the tip slightly penetrates the surface) may also contribute to the contrast in the height and phase images and this phenomenon may be at the origin of the increase of the apparent diameter of the hard spherical domains in the polymer blend thin films. For phase images recorded in the dominant repulsive regime, the phase shifts are related to the local mechanical properties. Below, we describe how to address more accurately this problem with PFQNM by imaging the stiffness, the adhesion, the deformation and the dissipation properties along with the topography of the polymer thin films in real-time. 3.3. Film properties of dry uncured coatings
Fig. 3. Height and phase TM-AFM images of thin films for the soft polymer (a) and (b) and a series of three blends with hard:soft ratios 20:80 (c) and (d), 30:70 (e) and (f), and 40:60 (g) and (h). The scan size is 2.0 m and the vertical scale for height images are 4.5 nm, 20 nm, 50 nm and 80 nm, respectively. For the phase images, the vertical scales are 9.0◦ for image (b) and 20◦ for (d), (f) and (h).
3.2. Film microstructure AFM [22–24] has been used for the elucidation of the phaseseparated microstructure in thin films of polymer blends. Tapping mode images can be of two different kinds: in the first one, the image corresponds to the changes of the piezoactuator height that are necessary to maintain a fixed oscillation amplitude, through a feedback loop (the height image); in the second one, the image contains the changes of the oscillator phase delay relative to the excitation signal (the phase image). In most cases, the phase measurement yields images reflecting tiny variations of the local (mechanical) properties of the sample surface. On that basis, it is possible to extract useful information from tapping mode phase images of soft samples, especially for samples showing compositional heterogeneity at small scale, for instance blends of hard and soft materials. Fig. 3a–h illustrates the typical results obtained on the soft polymer and on three systems with increasing hard:soft ratios (20:80, 30:70 and 40:60). The polymer films were prepared with a drying
Fig. 4a shows the real part of the complex heat capacity for the free-standing films of the references and mixed resins dried at 40 ◦ C. The thermograms show a single glass transition at Tg = 1 and 116 ◦ C for the soft and hard acrylated polyurethanes, respectively. In contrast, two glass transitions are identified for the blends and the weakness of the transitions lends support to the existence of a composition gradient forming an interphase between the soft continuum and harder domains. The Tg values of the blends are plotted as a function of composition in Fig. 4b. The lowest glass transition is attributed to a substantial fraction of soft polymer as suggested by the Tg -value, i.e. Tg,s ≈ 0–5 ◦ C, and from the amplitude of the heat capacity change. The glass transition temperature of the weaker transition at higher temperatures increases slightly from Tg,h ≈ 80 to 90 ◦ C for the films dried at 40 ◦ C. It is assigned to a phase primarily composed of hard polymer but to some extent plasticized by the soft polymer. The films dried at 80 ◦ C show a more pronounced phase contrast as inferred from the higher Tg of the hard phase (not shown). Here, Tg,h increases from 95 to 107 ◦ C upon increase of the hard:soft ratio from 20:80 to 40:60, closely approaching the Tg value of the pure hard polymer. The indentation profiles obtained by TMA in Fig. 5 show multiple softening steps. The profiles are essentially characterized by a prevailing indentation step but also weaker steps can be discerned. Three steps can be resolved for the blends whereas two steps are determined for the base polymers. Multiple indentation steps likely reflect the heterogeneous nature of the material. For the softer coatings, the first step is clearly related to the onset of softening of the
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Fig. 4. (a) Temperature dependence of the real part of the complex heat capacity for coatings prepared at a drying temperature of 40 ◦ C. The thermograms are shifted downwards with increasing polymer content. The blend composition is indicated by the hard:soft ratio. (b) Tg -values plotted as a function of hard:soft ratio. The low and high glass transition temperatures Tg,s and Tg,h are shown, respectively, by diamonds and triangles for the coatings dried at 40 ◦ C (solid) and 80 ◦ C (open symbols).
low-Tg (soft) compound. For the harder materials, however, the first weak dimensional change does not necessarily reflect indentation but is partly due to contact leveling before actual indentation. This explains the apparent onset of indentation for the hard polymer at low temperatures. The enrichment of the coating with hard polymer appears as an overall shift of the indentation curve towards higher temperatures and a decrease of the indentation in the high temperature plateau. As shown in Fig. 6a, the onset temperature of the major indentation step (softening temperature) ranges from 67 to 103 ◦ C across the investigated blend range which indicates that all the coatings behave essentially as hard materials at room temperature. The high softening temperature also suggests that the mechanical performance is ruled by the filling effect of the hard resin and likely the creation of a mixed interphase. Coatings dried at 80 ◦ C show a higher softening temperature for the blends. The onset temperature of the third indentation step suggests the presence of a hard phase in agreement with the DSC results. Interestingly, the plateau indentation at the high temperature end of the TMA curves exhibits a significant reduction up to a
hard:soft ratio ≈ 40:60 where the indentation hardness reaches a constant level (Fig. 6b). From this point on, there is no additional reinforcement of the coating due to the hard component which, from a maximum packing perspective, relates to the loss of connectivity in the soft continuum. It is noticed that the reinforcement is much steeper and nearly linear with respect to the mass fraction of the hard UV-PUD for coatings dried at 80 ◦ C. The proportional response of the hardness for coatings with well separated phase morphology is reminiscent of a filling reinforcement mechanism. As further discussed below, however, the marked hardness buildup at low hard:soft ratios cannot be explained by an inert filling effect alone. From the AFM and DSC data, it is argued that the formation of a composition gradient between the soft matrix and hard inclusions is essential to account for the hardness of these coating blends. The presence of such composition gradient would be consistent with the fact that the hard spheres seen with TM-AFM show an increasing diameter with the content of the hard component. The reinforcing effect resulting from the incorporation of the hard polymer is also established from the tensile properties shown in Fig. 7a. Starting from the soft polymer with a viscous-like behavior, a significant stress enhancement and a strong reduction of the ultimate elongation are observed with increasing hard polymer loading. As before, this trend is amplified for the coatings dried at 80 ◦ C. This is particularly well illustrated by the Young’s modulus in Fig. 7b, where the reinforcement results in a nearly straight increase for the coatings prepared at 80 ◦ C. The enhancement is less marked for coatings dried at 40 ◦ C, at least at low hard polymer content. Many phenomenological and micromechanical models have been suggested to describe the relation between composition and tensile modulus for immiscible polymer blends and filled polymers [7–11]. This area has been reviewed extensively. A simple upper bound is provided by the parallel or rule-of-mixture (ROM) model predicts a linear relationship for the modulus of the composite or blend material, i.e., Ec,ROM = m Em + f Ef
Fig. 5. TMA indentation profiles of free-standing coatings prepared from the base polyurethanes and their blends. The coatings dried at 40 and 80 ◦ C are shown by solid and dashed lines, respectively. The thickness of the lines indicates increasing contents of hard materials, i.e. 0:100, 20:80, 30:70, 40:60. The hard polyurethane is shown by a double line.
(1)
where m and f = 1 − m stand for the volume fractions of the matrix continuum and the dispersed (filling) compound, respectively, Em and Ef are the tensile moduli of the matrix and the dispersed material, respectively. In this particular case, it is assumed that the two phases deform identically with equal strain. In the opposite limiting situation (lower bound), the two material
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Fig. 6. (a) Composition dependence of the softening temperature. The onset temperature of the major indentation step is shown with large symbols whereas the smaller symbols (diamonds, triangles) are used for the weaker steps. Coatings dried at 40 and 80 ◦ C are represented by solid and open symbols, respectively. (b) High temperature indentation plateau as a function of hard:soft polymer ratio. Coatings dried at 40 and 80 ◦ C are shown by diamonds and triangles, respectively.
components experience equal stress and the composite modulus is then given by Ec,IROM =
Em Ef m Ef + f Em
(2)
This model, known as the inverse rule-of-mixture (IROM), is actually the series equivalent of the parallel model. Fig. 7b shows that the experimental data are contained between the two limits provided by Eqs. (1) and (2). For further comparison, the semiempirical Halpin–Tsai (HT) expression [8] for an isotropic material of spherical particles dispersed in a matrix is also plotted (dashed line). The composite HT modulus reads as Ec,HT = Em
Ef + (m Em + f Ef ) ( + f )Em + m Ef
(3)
where is a parameter related to the geometry and the packing of the filling compound. In Fig. 7b, the dispersed particles were assumed as spherical with a value = 2. The HT equation reduces to Eq. (1) for → ∞ and to Eq. (2) for → 0. Leaving as an adjustable parameter in HT equation does not result in a good fit of the experimental data. The reinforcement effect due to the hard component
grows faster with polymer mass fraction than expected from the Halpsin–Tsai and IROM predictions when volume and mass fractions are considered as equal. This possibly suggests that the hard polymer mass fraction underestimates the effective volume fraction of the dispersed component in the coating. The presence of a compositional gradient (interphase) could partly account for the higher filling volume. However, diffuse phase boundaries are usually beyond the assumptions of the traditional micromechanical models. A notable exception is the self-consistent mechanical model elaborated by Colombini and Maurer which accounts for the occurrence of so-called micromechanical viscoelastic transitions by invoking an interlayer in multiphase polymer materials [25]. Simulations show that an interlayer with sufficient strength can indeed lead to reinforcement but not to the extent that is observed here. Following a simple approach, we have modified the twospring matrix-filler model, represented by Eqs. (1) and (2), into a three compartment model with an additional elastic spring for the interphase region. The experimental data is most conveniently described by a parallel combination of the low (matrix) and medium modulus compartments (interphase) in series with
Fig. 7. (a) Tensile properties of the free-standing films at room temperature. The coatings dried at 40 and 80 ◦ C are shown by solid and dashed lines, respectively. An increasing line thickness refers to an increasing hard:soft ratio. (b) Plot of Young’s Modulus as a function of the hard polymer content in the coating. Drying temperature 40 ◦ C (circles), 80 ◦ C (triangles). The experimental data are compared to the following empirical predictions: ROM (upper dotted line), IROM (lower dotted line), Halpin–Tsai (HT: thick dashed line). The solid lines through the data result from a non-linear least-squares comparison to the three-compartment (TC) model described by Eq. (4).
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Fig. 8. Illustration of some parameters which can be extracted in real time for each tapping cycle by using the PFQNM technique.
the high modulus compartment (filler). The resulting composite modulus reads as f m + int 1 = + Ef Ec Em,int
(4)
where Em,int = (m Em + int Eint )/(m + int ) is the modulus value of the interphase and the matrix in parallel combination. Young’s modulus of the matrix and the filler are, respectively, Em ≈ 0.05 GPa (measured) and Ef ≈ 2 GPa (typical value for a glassy UVPUD). It is assumed that the effective volume fractions f and int are proportional to the mass fraction wf of the dispersed hard polymer, i.e., f = Ff wf , int = Fint wf and m = 1 − f − int . The full lines through the data in Fig. 7b result from a least-squares comparison to Eq. (4) using the proportionality factors Ff and Fint as adjustable parameters. An arithmetic mean value Eint ≈ 1 GPa is tentatively taken for the average modulus of the compartment in-between the filler and matrix compartments. The enhancement of Young’s modulus at low hard polymer content is well captured by this phenomenological model. When comparing the values of the fit parameters at 40 and 80 ◦ C, it appears that the factor Ff reduces from 1.73 to 0.86 whereas Fint inflates from 0.11 to 0.77. These values suggest that the expansion of the interphase at 80 ◦ C is the reason behind the significant enhancement of Young’s modulus compared to the films at 40 ◦ C. The reduction of the filler fraction Ff mainly shifts the divergence to higher mass fractions and has little effect on the behavior at low mass fractions. AFM imaging with Peak Force Quantitative Nanomechanical Property Mapping (PFQNM) [26] is based on the real-time analysis of the force–distance curves recorded at a frequency of about 2 kHz (while usually these curves are recorded at 1 Hz). Since the feedback loop is set on the (peak) force, the actual force is maintained constant during the imaging of the samples. From those force–distance curves, one can extract the following information: (i) the adhesion force (corresponding to the lowest value of the force when the tip is retracted from the sample; (ii) Young’s modulus (obtained from the slope of the curve while the tip is “indenting” the sample, according to the DMT model [27]); (iii) the dissipation (corresponding to the surface integral between the approach and the retract curves); and (iv) the deformation (by comparing the curves to those obtained for a very stiff sample). These parameters are illustrated on Fig. 8. It is important to note that other methods based on a multi-frequency signal sent to the cantilever are also available like Dual AC [28,29], Intermodulation [30] or Band Excitation [13] but it is extremely difficult to extract quantitative values out of the data even if they are providing a contrast in the images. Since the force–distance curves are recorded at every point of the raster scan, the analysis provides a map of those properties. As an example, Fig. 9 illustrates the data obtained for the coating with
a hard:soft ratio of 20:80 at two drying temperatures (left: 40 ◦ C; right: 80 ◦ C). The scale range is 20 nm for the two height images. It is important to note that the topographic images obtained by PFQNM are equivalent to those recorded using TM-AFM. On the adhesion images of the blend with a hard:soft ratio of 20:80, the hard spherical particles are much less adherent than the soft polymer matrix and therefore appear in black in the image. By comparing the roughness of the samples dried at 40 ◦ C (3.25 nm) to the value measured for a drying temperature of 80 ◦ C (1.75 nm), we can conclude that the drying temperature drastically influences the flatness of the dried films. This observation is independent from the hard:soft ratio. More interestingly, the adhesion image for 80 ◦ C (Fig. 9d, right) is less contrasted compared to that of the sample dried at 40 ◦ C and the apparent diameter of the spherical hard particles is smaller. This indicates that the hard spheres are more embedded in the soft film, which is also the main reason why the DMT modulus image is less contrasted. This is fully consistent with the model proposed above, which considers the presence of an interphase between the hard spheres and the soft polymer matrix when the drying temperature is 80 ◦ C. This interpretation is also in agreement with the data from Colombini et al. where they carefully studied the morphology and the viscoelastic properties of bimodal hard/soft latex blends [21]. The latter blends can be considered here as a model compound relative to our UV-PUD systems. 3.4. Film properties of dry UV-cured coatings Regardless of the composition and the drying temperature, all the coatings exhibited a similar thermal behavior after UV-curing characterized by a single glass transition with a Tg value in the range 106–125 ◦ C. The initially soft UV-PUD undergoes a remarkable Tg -increase of more than 100 ◦ C upon UV-curing. Hence, the high crosslink density in both soft and hard regions levels off any compositional contrast, leaving behind an apparently homogenous and extremely hard coating which prevents any indentation experiment. Finally, the tensile behavior is overall hard and brittle with Young’s modulus values of ≈2.2–2.6 GPa and ultimate elongations down to approximately 1%. 3.5. Film performance The soft polymer was compared with the 40:60 blend which was selected as the most promising composition from this study. The polymer dispersions were formulated into orange pigmented compositions for spray application, the choice of the orange color being guided by the fact that it is a challenging color for UVpolymerization owing to the strong competition between the UV-absorption of the photoinitiator and the pigments. For that purpose, they were mixed with an industrial ready-to-use water-based orange pigment paste (15%, w/w) and further formulated with two liquid photoinitiators (1% ESACURETM HB and 0.5% IRGACURETM 819 DW), a rheology modifier (0.5% UCECOATTM 8460) and a wetting agent (1% BYKTM 028). The pigmented formulations were applied using a spray gun technique to MDF panels laminated with white melamine paper and prepared by sanding with a thin aluminum oxide abrasive paper. The coating with a wet thickness of 120–130 g m−2 was dried in a ventilated oven for 20 min at 40 ◦ C and immediately cured on a conveyor belt at a speed of 5 m min−1 using a Ga-doped mediumpressure mercury lamp (power input: 120 W cm−1 ) followed by a medium-pressure mercury lamp (power input: 120 W cm−1 ). The cured coatings were tested for the main performance areas, i.e. adhesion (DIN 53151), stain resistance (EN12720) and nail scratch resistance. The results of the tests are reported on a 0–5 scale (5 = excellent). The adhesion and the stain resistance (black marker, red wine and coffee) of the two colored coatings were all
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Fig. 9. Left: PFQNM images of a thin film of the polyurethane blend with a hard:soft ratio of 20:80. The film was dried at 40 ◦ C. (a) Height; (b) adhesion; (c) DMT modulus; and (d) dissipation. The scan size is 1.0 m. Right: PFQNM images of a thin film of the polyurethane blend with a hard:soft ratio of 20:80. The film was dried at 80 ◦ C. (a) Height; (b) adhesion; (c) DMT modulus; and (d) dissipation. The scan size is 1.0 m. Brighter areas correspond to higher relative values of the image property.
It is clear that the resulting product can be positioned at the high-end performance level for use in pigmented coatings, for instance on wood furniture. It exhibits an excellent colloidal stability and a low viscosity so that it can be easily applied by spray or curtain coating. It is non-irritant and does not contain any cosolvent. It has a low film formation temperature and is tack-free before cure. Importantly, it does not require the use of additional expensive thermal crosslinkers detrimental to the pot-life of the formulation and to occupational health and safety. Acknowledgments
Fig. 10. Comparison of the nail scratch resistance of orange UV-cured coatings based on the soft polymer (0:100) and the blend with the hard:soft ratio of 40:60. The results are plotted as a function of elapsed time (0–5 scale, 5 = best).
found to be excellent (5/5). The results of the nail scratch resistance are shown in Fig. 10 and address the through cure efficiency in the presence of pigments. The performance of the two products increases over the investigated time interval but the blend gives significantly better results immediately after the curing (when the panel is still hot) and reaches a maximum value after 60 min, allowing a good stackability of the coated panels during their manufacture. 4. Conclusions We developed new aliphatic radiation curable polyurethane dispersions by combining soft and hard particles in a blend composition. The blends showed a low film formation temperature associated with a very significant enhancement of the mechanical properties before cure for a hard:soft ratio lower than 50:50. This reinforcement can be explained by the formation of a microstructure with an interphase around the hard spherical particles and is supported by PFQNM AFM data. An optimal ratio was found between the soft and hard components showing at the same time an exceptional chemical and mechanical resistance after curing in the presence of difficult pigments.
The authors wish to thank all the collaborators who were involved with syntheses, application & testing as well as all of the colleagues who gave us their trust and support to develop this new technology. In particular, we are thankful to G. Hellin who made his graduation work both in the research laboratory at CYTEC Surface Specialties and in the laboratory at the University of Mons. We are also grateful to Dr. Chanmin Su (Bruker Nano Inc., Santa Barbara, CA) for helpful discussions on PFQNM. We specially thank Formulaction for their great help in the investigation using ASII. Ph. Leclère is FRS-FNRS Research Associate (Belgium). References [1] K. Buysens, M. Tielemans, Th. Randoux, Pitture e Vernici – European Coatings 19 (2002) 27. [2] M. Tielemans, P. Roose, Ph De Groote, J.-C. Vanovervelt, Prog. Org. Coat. 55 (2006) 128. [3] M. Tielemans, P. Roose, Prog. Org. Coat. 55 (2008) 182. [4] Ph. De Micheli, S. Peeters, Polymer Paint Colour J. 4 (2009). [5] M. Tielemans, J.-P. Bleus, M. Vasconi, Eur. Coat. J. 3 (2007) 38. [6] M. Tielemans, M. Decaux, Ph. De Micheli, S. Peeters, Eur. Coat. J. 4 (2009) 40. [7] R.A. Dickie, J. Appl. Polym. Sci. 17 (1973) 45. [8] J.C. Halpin, J.L. Kardos, Polym. Eng. Sci. 16 (1976) 344. [9] H. Eklind, F.H.J. Maurer, Polymer 37 (1996) 2641. [10] A.G. Facca, M.T. Kortschot, N. Yan, Composites A 37 (2006) 1660. [11] S.-Y. Fu, X.-Q. Feng, B. Lauke, Y.-W. Mai, Composites B 39 (2008) 933. [12] F. Benmouna, T.D. Dimitrova, D. Johannsmann, Langmuir 19 (2003) 10247. [13] S. Jesse, S.V. Kalinin, R. Proksch, A.P. Baddorf, B.J. Rodriguez, Nanotechnology 18 (2007) 435503. [14] R.M. Fitch, Polymer Colloids: A Comprehensive Introduction, Academic Press, London, 1997, pp. 314–337. [15] T. Goverde, U., Boetcher, M. Fleury (Formulaction), private communication.
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[16] H. Clausen-Schaumann, M. Seitz, R. Krautbauer, H.E. Gaub, Curr. Opin. Chem. Biol. 4 (2000) 524. [17] S. Ge, Y. Pu, W. Zhang, M. Rafaidovich, J. Sokolov, C. Buenviaje, R. Buckmaster, R.M. Overney, Phys. Rev. Lett. 85 (2000) 2340. [18] T. Kajiyama, K. Tanaka, I. Ohki, S.R. Ge, J.S. Yoon, A. Takahara, Macromolecules 27 (1994) 7932. [19] M. Radmacher, R.W. Tillmann, H.E. Gaub, Biophys. J. 64 (1993) 735. [20] J. Mallegol, O. Dupont, J.L. Keddie, Langmuir 17 (2001) 7022. [21] D. Colombini, H. Hassander, O.J. Karlsson, F.H.J. Maurer, Macromolecules 37 (2004) 6865. [22] S. Kopp-Marsaudon, Ph Leclère, F. Dubourg, R. Lazzaroni, J.P. Aimé, Langmuir 16 (2000) 8432.
[23] Ph. Leclère, F. Dubourg, S. Kopp-Marsaudon, J.L. Brédas, R. Lazzaroni, J.P. Aimé, Appl. Surf. Sci. 188 (2002) 524. [24] Ph. Leclère, P. Viville, M. Jeusette, J.P. Aimé, R. Lazzaroni, Scanning Probe Microscopies: Beyond Imaging, Wiley-VCH, 2006, pp. 175–207. [25] D. Colombini, F.H.J. Maurer, Macromolecules 35 (2002) 5891. [26] http://www.bruker-axs.com/uploads/tx linkselectorforpdfpool/PeakForce Quantitative Nanomechanical Property Mapping brochure.pdf. [27] B.V. Derjaguin, V.M. Muller, Y.P. Toropov, J. Colloid. Interface Sci. 53 (1975) 314. [28] B.J. Rodriguez, C. Callahan, S.V. Kalinin, R. Proksch, Nanotechnology 18 (2007) 475504. [29] J.R. Lozano, R. Garcia, Phys. Rev. Lett. 100 (2008) 076102. [30] D. Platz, E.A. Tholén, D. Pesen, D.B. Haviland, Appl. Phys. Lett. 92 (2008) 153106.
Contents lists available at SciVerse ScienceDirect
Progress in Organic Coatings journal homepage: www.elsevier.com/locate/porgcoat
Multiphase coatings from complex radiation curable polyurethane dispersions Michel Tielemans a,∗ , Patrice Roose a , Chinh Ngo b , Roberto Lazzaroni b , Philippe Leclère b,∗∗ a b
R&D, Cytec Surface Specialties, Anderlechtstraat 33, B-1620 Drogenbos, Belgium Laboratory for Chemistry of Novel Materials, University of Mons, Place du Parc 20, B-7000 Mons, Belgium
a r t i c l e
i n f o
Article history: Received 10 August 2011 Received in revised form 25 April 2012 Accepted 22 May 2012 Available online 15 June 2012 Keywords: Polyurethane Waterborne Radiation-curing Morphology Film-formation Atomic force microscopy Pigments
a b s t r a c t The waterborne nature of radiation curable polyurethane dispersions largely respond to the current environmental concerns and do not require any additional coalescent since the film formation (drying) and hardening (photo-curing) take place in distinct steps. It is possible to design aqueous dispersions with distinct polymer particle populations resulting in micro-structured coatings with optimized properties over a wide range of curing conditions. Mixed dispersions based on hard and soft acrylated polyurethane particles were used as model systems for the present study. The minimum film formation temperature has been investigated as a function of the hard:soft polymer ratio. The elastic modulus of the dry coatings shows a reinforcing effect consistent with the inclusion of hard domains in a soft continuous matrix. However, the level of reinforcement is not properly predicted by the usual mechanical models and it is qualitatively accounted for by assuming a composition gradient (interphase) between the hard domains and the matrix. The multiple-phase structure was clearly established by Atomic Force Microscopy in agreement with thermal analysis data. Furthermore the local nanoscale mechanical properties were mapped using a new imaging mode based on real-time force–distance curve analysis. Finally, the coatings prepared using this multiple-phase pattern present a clear benefit over conventional homogeneous coatings by offering an improved balance of chemical and mechanical resistance in pigmented systems applied on melamine-coated MDF panels. © 2012 Elsevier B.V. All rights reserved.
1. Introduction The introduction of water-based alternatives profoundly changed the landscape of the radiation curing market by combining the high productivity and performance of traditional radiation curing compositions with the low viscosity of usual water-based systems [1]. Aqueous colloidal dispersions of acrylated polyurethane oligomers (referred to as UV-PUD’s) largely respond to the environmental regulations due to their waterborne nature and the absence of volatile organic compounds or additional coalescents, since the film formation (drying) and hardening (photo-curing) take place in distinct steps. They have a good colloidal stability [2] and their rheology [3] makes them particularly suitable for application by spray, curtain or roller onto different substrates. A wide range of products has been developed for clearcoat application on wood furniture with excellent adhesion, chemical and mechanical resistance for indoor [4] as well as outdoor [5] applications.
∗ Corresponding author. ∗∗ Corresponding author. E-mail addresses: [email protected] (M. Tielemans), [email protected] (P. Leclère). 0300-9440/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.porgcoat.2012.05.010
The coatings obtained after film formation and photocuring are characterized by exceptional mechanical properties since they combine the usual features of polyurethanes along with high cross-linking levels. The microstructure of the cured coating is typically characterized by an alternation of hard and soft segments entrapped within a dense crosslinked network, providing a particular ability to balance contrasting properties, such as flexibility and adhesion, to various substrates on one side, scratch resistance and chemical resistance on the other side. As long as clearcoats are concerned, light penetration and in-depth curing (“through cure”) of the coating is not an issue. However, a technological limitation arises when the radiation curable polyurethane dispersions are formulated with organic or inorganic pigments so as to obtain a colored decorative coating. In this case, substantial ultraviolet light is absorbed and scattered by the pigments which prevents efficient light penetration into the coating, in particular for pigments other than white. In order to design a colored coating with sufficient mechanical strength in the bulk along with strong chemical resistance at the surface, without sacrificing the film formation and the adhesion at the coating–substrate interface, a mixing approach with tuned particle characteristics has been suggested recently [6]. It is possible to make from soft to very hard polymer particles dispersed in water and presenting each the characteristics inherent to their chemical nature. Soft particles will provide low
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other existing and promising AFM techniques such as AFM Noise Analysis [12] or Band Excitation AFM [13]. Finally, the product has been formulated with a high pigment load and applied by spray on melamine-coated MDF panels prior to the curing under UV light. The coatings prepared using a multiplephase structure of hard and soft acrylated polymeric components present a distinctive benefit over conventional homogeneous coatings. 2. Materials and methods
Fig. 1. Schematic representation of polymer dispersions made from (a) soft particles, (b) hard particles, (c) hard core–soft shell particles and (d) soft and hard particles.
film formation temperatures without any coalescing solvent but the mechanical properties of the bulk film before cure (or in the absence of cure in case of strong light interaction with pigments) are very poor. In contrast, hard particles will enhance the mechanical properties before cure but the high film formation temperature, requesting significant levels of coalescent, is totally inappropriate and usually associated with a poor adhesion and depressed optical properties. It appears that only complex polymer dispersion morphologies are suited to balance the coating properties in the appropriate manner. Although core-shell morphologies offer some real attractiveness for designing multiple-phase coatings, they are much more difficult to control and to manufacture so that their benefits over particle blending are not always straightforward. The dispersions originating from soft particles, hard particles, hard core–soft shell particles and a physical blend of soft and hard particles are schematically depicted in Fig. 1. The UV-PUD system used in the present study is based on a blend of hard (i.e., with a glass transition temperature, Tg , higher than room temperature) and soft (with Tg lower than room temperature) acrylated aliphatic polyurethane particles and demonstrates that well balanced properties can be achieved before and after UV-curing following this approach. In this paper, we first examine the minimum film formation temperature (MFFT) as a function of the hard:soft ratio in the blend composition. Next, in order to characterize the phase structure of the coatings, the thermal and thermo-mechanical properties are measured by temperature-modulated differential scanning calorimetry (TMDSC) and thermo-mechanical analysis (TMA). The tensile properties show clearly a reinforcing effect of hard domains embedded in a continuous matrix. Many empirical or semi-empirical models have been proposed to predict the modulus of particulate–polymer composites [7–11]. However, the assumption of a composition gradient (“interphase”) between the hard inclusions and the soft matrix was required in order to capture the enhancement of Young’s modulus. The results have been correlated with the microphase characteristics determined by Atomic Force Microscopy (AFM). Moreover, the local mechanical properties, such as stiffness, adhesion, dissipation and deformation, were mapped by using a new imaging mode based on real-time force–distance curve analysis. This recent technique, referred to as Peak Force Quantitative Nanomechanical Property Mapping (PFQNM), returns local mechanical properties at the nanoscale within reasonable acquisition times compared to
UV-PUD’s are anionically stabilized acrylated polyurethane colloids in water [14]. A typical polyurethane ionomer structure is composed from (i) a polyacrylated molecule having at least one chemical functionality able to react with isocyanates; (ii) a polyisocyanate; (iii) a polyol selected among polyesters, polyethers or polycarbonates; (iv) a functional polyol capable to disperse the oligomer in water and whose functionality is usually a carboxylic acid or a salt thereof and (v) a polyamine (optional). They encompass a very large range of composition and molar mass with a linear or branched polymeric architecture and a varying amount of urethane, urea, allophanate and biuret functionality. The preparation of the urethane–acrylate polymer follows a multi-step process in a low boiling point solvent (e.g. acetone). The resulting polymer solution is then neutralized, dispersed and optionally chain-extended in water, the solvent being finally removed from the dispersion medium under reduced pressure. Two base dispersions were prepared for the purpose of this study, each originating from the same carboxylated and acrylated polyurethane backbone but clearly differentiated by their molar mass distributions. The first model dispersion is a soft (low Tg ) polymer dispersion with a low molar mass and a minimum film formation temperature (MFFT) below ambient temperature. The second model dispersion is a hard polymer dispersion with a very high molar mass range and a MFFT much higher than room temperature. The main characteristics of the two model dispersions are summarized in Table 1. The solid content of the dispersions was determined gravimetrically after 2 h at 120 ◦ C. The viscosity was measured with a Brookfield viscometer at 25 ◦ C and at a rotational speed of 50 rpm. The hydrodynamic particle diameter (dDLS ) was determined by dynamic light scattering using a Malvern Autosizer LoC instrument linked to a Malvern Series 7032 Multi-8 correlator. The colloidal stability was assessed by multiple light scattering (TurbiscanTM ) at 60 ◦ C by recording the number of days before critical destabilization. The weight-average molar mass (Mw ) was measured by gel permeation chromatography on the soluble polymer fraction in tetrahydrofuran. The two base dispersions have been investigated separately or mixed as blends in various ratios in order to characterize the film formation properties and to relate the microstructure to the mechanical properties in the dry state after film formation. A series of blends were prepared according to the hard:soft polymer ratio of 20:80, 30:70 and 40:60. When required, a photo-reactive
Table 1 Characteristics of the proprietary polyurethane dispersions used for the blends.
Solids (%, w/w) Viscosity (mPa·s) pH dDLS (nm) Stability at 60 ◦ C (days) Mw (kDa) Solubility in THF (%)
Soft
Hard
35 100 7.5 70 >10 15 100
35 100 8.0 60 >10 >100 ∼15
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formulation was obtained by adding 1.5% (w/w) ESACURETM HB as photoinitiator to the dispersions. The MFFT was measured according to ASTM 2354 on a Rhopoint® automatic gradient-heated metal plate. The film formation kinetics of the dispersions has been addressed using adaptive speckle imaging interferometry (ASII) by applying coatings (100 m) on glass at 22 ◦ C [15]. This technique is based on laser light scattering of a drying film where the resulting interference image is recorded with a camera detector. The interference pattern (speckle) changes with the reduction in mobility of the scattering elements and eventually provides a qualitative characterization of the film formation over time. Atomic Force Microscopy (AFM) is increasingly used for performing mechanical studies of materials on the nanometer scale. Widely known are single molecule force spectroscopy [16], shear force modulation [17], friction microscopy [18], force modulation [19] and slightly less quantitative, but providing pertinent information, the intermittent contact or Tapping Mode [20]. For the morphological characterization of the film microstructure, Tapping-Mode AFM (TM-AFM) was used on a Veeco (now Bruker Nano Inc.) ICON AFM (Santa Barbara, CA) driven by a Nanoscope V control unit. We operated in air and in ambient conditions (temperature and pressure). Thin coating films were prepared on cleaned glass substrates using a bar coater and subsequently dried at 40 or 80 ◦ C. The probe has a resonance frequency of about 300 kHz. At this point it is worth mentioning that, in order to keep a well-defined oscillating behavior of the tip, the perturbation to the oscillator due to the contact with the surface is chosen to be small. In other words, the reduction with respect to the free amplitude (the setpoint) is only of a few percent (typically 5%). This method has two advantages. From an experimental point of view, this allows to identify easily hard and soft areas, the strongly phase-shifted areas of the image corresponding to the hard domains. From a theoretical point of view, it permits to use simple approximations providing analytical solutions that help to interpret the experimental data. Free-standing coating films were prepared for the determination of the thermo-mechanical properties of the different systems. The aqueous UV-PUD resins were applied on an untreated bioriented polypropylene (BOPP) plastic film (30 m) using a bar coater. In order to obtain films of suitable thickness, five layers of ≈35 m (wet) were applied on top of each other and every layer was dried for 1 h in an air-convection oven. The effect of the drying temperature was studied at 40 and 80 ◦ C. The liquid polymer dispersion showed good wetting on the BOPP film and the coatings could easily be removed from the substrates after drying. A portion of every film was separated and cured twice on a UV-conveyor belt using a high power medium-pressure mercury lamp (power input 80 W cm−1 ; belt speed 5 m min−1 ). The phase behavior of the dried coatings was studied using temperature-modulated differential scanning calorimetry (TMDSC). The measurements were conducted using a Mettler DSC823e instrument with standard aluminum crucibles and nitrogen as the flow gas (50 mL min−1 ). The sample mass was typically ≈10 mg. The sample was heated using a sinusoidal temperature modulation with a period of 60 s and an amplitude of 1 ◦ C superimposed on a linear ramp of 2 ◦ C min−1 . Prior to the modulated heating step, the sample was heated at 85 ◦ C for 5 min in order to improve samplepan contact. The glass transition temperature is determined at the midpoint of the transition. Thermo-mechanical analysis (TMA) was conducted in static penetration mode using a flat-ended indentation rod (radius 1.27 mm) under a static normal load of 0.5 N. TMA data were recorded with a 2940 thermo-mechanical analyzer (TA instruments) at a heating rate of 3 ◦ C min−1 . Tensile properties were measured at room temperature with a Zwick Z010 elongation testing machine at a cross-head speed
of 1 mm min−1 . For the tests, 3 cm × 0.5 cm rectangular-shaped specimens were cut from the coatings. At least five independent measurements were performed for each sample.
3. Results and discussion 3.1. Film formation The dispersions prepared from the soft and hard acrylated polyurethanes were investigated for MFFT. The soft polymer dispersion has a MFFT value lower than 0 ◦ C as common for this kind of composition. On the other hand, the MFFT value of the hard polymer dispersion is beyond 90 ◦ C. The basic idea behind our approach is that the non film-forming component can be incorporated through the presence of the film-forming component. The MFFT has been investigated as a function of the hard:soft polymer ratio. The results are plotted in Fig. 2. The data present a gradual but slow increase of the MFFT up to ≈50% (w/w) of hard polymer (i.e., hard:soft ratio = 50:50) and then exhibit a marked rise after the transition point. This observation is consistent with a soft polymer matrix surrounding the hard particles up to the critical volume concentration where there is not enough of the soft polymer left to fill effectively the space between the hard polymer particles. Colombini et al. showed that the particle size ratio of hard and soft latex blends also has a large influence on the MFFT behavior [21]. They found that for a blend with a 50:50 hard:soft polymer ratio, the MFFT does not depart very much from the value of the soft matrix as long as the particle size ratio dsoft /dhard stays below 1.5. In agreement, the MFFT of the 50:50 blend in Fig. 2 has raised by 12 ◦ C compared to the pure soft polymer but still suggests that the continuum is essentially formed by the soft compound. Beyond this point, the formation of hard particle aggregates and later percolation leads to bicontinuous and inverse morphologies with a marked enhancement of the MFFT as a consequence [21]. For all the compositions further analyzed in this study, the soft polymer is thus effectively embedding the hard polymer by creating a continuous film forming matrix with hard polymer inclusions as observed by AFM (vide infra). The ASII experiments conducted for the soft polymer dispersion and the blend with a hard:soft ratio of 40:60 support this statement since no significant differences could be recorded from the shape of the drying profiles, showing similar characteristic drying times at 22 ◦ C.
Fig. 2. Minimum film formation temperature of coating blends as a function of hard polyurethane content (%, w/w).
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temperature of 40 ◦ C. The thin film of the pure soft polymer is flat with small, shallow holes; the surface is characterized by a rootmean-square roughness of only 0.35 nm. The roughness increases to 3.25 nm and 8.35 nm when incorporating the hard polymer at 20% (w/w) and 40% (w/w), respectively. The morphology observed on the height image (c), (e) and (g) corresponds to a dispersion of small spherical objects appearing harder than the surrounding matrix on the corresponding phase images (d), (f) and (h). From the image analysis, we can also conclude that the number of these spherical objects increases with the volume fraction of hard polymer present in the blend. For the 40:60 ratio (images (g) and (h)), the spheres are packed very close to each other. The apparent average diameter of the spherical objects is also slightly increasing with the hard:soft ratio (56 nm for 20:80, 58 nm for 30:70 and 60 nm for 40:60). The samples dried at 80 ◦ C (images not shown here) are very similar in terms of morphology and average diameter of the hard domains (58 nm for 20:80, 62 nm for 30:70 and 65 nm for 40:60). From these results, we can conclude that the hard polymer is well micro-phase separated from the soft polymer matrix and forms spherical objects visible in TM-AFM images with increasing apparent diameters relative to the hard:soft ratio. Despite the success of TM-AFM for micro-structural characterization, important questions remain about the physical origin of the image contrast. The height images are generally considered to display topographic information, but it must be kept in mind that the local mechanical properties of the sample (i.e., the possibility that the tip slightly penetrates the surface) may also contribute to the contrast in the height and phase images and this phenomenon may be at the origin of the increase of the apparent diameter of the hard spherical domains in the polymer blend thin films. For phase images recorded in the dominant repulsive regime, the phase shifts are related to the local mechanical properties. Below, we describe how to address more accurately this problem with PFQNM by imaging the stiffness, the adhesion, the deformation and the dissipation properties along with the topography of the polymer thin films in real-time. 3.3. Film properties of dry uncured coatings
Fig. 3. Height and phase TM-AFM images of thin films for the soft polymer (a) and (b) and a series of three blends with hard:soft ratios 20:80 (c) and (d), 30:70 (e) and (f), and 40:60 (g) and (h). The scan size is 2.0 m and the vertical scale for height images are 4.5 nm, 20 nm, 50 nm and 80 nm, respectively. For the phase images, the vertical scales are 9.0◦ for image (b) and 20◦ for (d), (f) and (h).
3.2. Film microstructure AFM [22–24] has been used for the elucidation of the phaseseparated microstructure in thin films of polymer blends. Tapping mode images can be of two different kinds: in the first one, the image corresponds to the changes of the piezoactuator height that are necessary to maintain a fixed oscillation amplitude, through a feedback loop (the height image); in the second one, the image contains the changes of the oscillator phase delay relative to the excitation signal (the phase image). In most cases, the phase measurement yields images reflecting tiny variations of the local (mechanical) properties of the sample surface. On that basis, it is possible to extract useful information from tapping mode phase images of soft samples, especially for samples showing compositional heterogeneity at small scale, for instance blends of hard and soft materials. Fig. 3a–h illustrates the typical results obtained on the soft polymer and on three systems with increasing hard:soft ratios (20:80, 30:70 and 40:60). The polymer films were prepared with a drying
Fig. 4a shows the real part of the complex heat capacity for the free-standing films of the references and mixed resins dried at 40 ◦ C. The thermograms show a single glass transition at Tg = 1 and 116 ◦ C for the soft and hard acrylated polyurethanes, respectively. In contrast, two glass transitions are identified for the blends and the weakness of the transitions lends support to the existence of a composition gradient forming an interphase between the soft continuum and harder domains. The Tg values of the blends are plotted as a function of composition in Fig. 4b. The lowest glass transition is attributed to a substantial fraction of soft polymer as suggested by the Tg -value, i.e. Tg,s ≈ 0–5 ◦ C, and from the amplitude of the heat capacity change. The glass transition temperature of the weaker transition at higher temperatures increases slightly from Tg,h ≈ 80 to 90 ◦ C for the films dried at 40 ◦ C. It is assigned to a phase primarily composed of hard polymer but to some extent plasticized by the soft polymer. The films dried at 80 ◦ C show a more pronounced phase contrast as inferred from the higher Tg of the hard phase (not shown). Here, Tg,h increases from 95 to 107 ◦ C upon increase of the hard:soft ratio from 20:80 to 40:60, closely approaching the Tg value of the pure hard polymer. The indentation profiles obtained by TMA in Fig. 5 show multiple softening steps. The profiles are essentially characterized by a prevailing indentation step but also weaker steps can be discerned. Three steps can be resolved for the blends whereas two steps are determined for the base polymers. Multiple indentation steps likely reflect the heterogeneous nature of the material. For the softer coatings, the first step is clearly related to the onset of softening of the
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Fig. 4. (a) Temperature dependence of the real part of the complex heat capacity for coatings prepared at a drying temperature of 40 ◦ C. The thermograms are shifted downwards with increasing polymer content. The blend composition is indicated by the hard:soft ratio. (b) Tg -values plotted as a function of hard:soft ratio. The low and high glass transition temperatures Tg,s and Tg,h are shown, respectively, by diamonds and triangles for the coatings dried at 40 ◦ C (solid) and 80 ◦ C (open symbols).
low-Tg (soft) compound. For the harder materials, however, the first weak dimensional change does not necessarily reflect indentation but is partly due to contact leveling before actual indentation. This explains the apparent onset of indentation for the hard polymer at low temperatures. The enrichment of the coating with hard polymer appears as an overall shift of the indentation curve towards higher temperatures and a decrease of the indentation in the high temperature plateau. As shown in Fig. 6a, the onset temperature of the major indentation step (softening temperature) ranges from 67 to 103 ◦ C across the investigated blend range which indicates that all the coatings behave essentially as hard materials at room temperature. The high softening temperature also suggests that the mechanical performance is ruled by the filling effect of the hard resin and likely the creation of a mixed interphase. Coatings dried at 80 ◦ C show a higher softening temperature for the blends. The onset temperature of the third indentation step suggests the presence of a hard phase in agreement with the DSC results. Interestingly, the plateau indentation at the high temperature end of the TMA curves exhibits a significant reduction up to a
hard:soft ratio ≈ 40:60 where the indentation hardness reaches a constant level (Fig. 6b). From this point on, there is no additional reinforcement of the coating due to the hard component which, from a maximum packing perspective, relates to the loss of connectivity in the soft continuum. It is noticed that the reinforcement is much steeper and nearly linear with respect to the mass fraction of the hard UV-PUD for coatings dried at 80 ◦ C. The proportional response of the hardness for coatings with well separated phase morphology is reminiscent of a filling reinforcement mechanism. As further discussed below, however, the marked hardness buildup at low hard:soft ratios cannot be explained by an inert filling effect alone. From the AFM and DSC data, it is argued that the formation of a composition gradient between the soft matrix and hard inclusions is essential to account for the hardness of these coating blends. The presence of such composition gradient would be consistent with the fact that the hard spheres seen with TM-AFM show an increasing diameter with the content of the hard component. The reinforcing effect resulting from the incorporation of the hard polymer is also established from the tensile properties shown in Fig. 7a. Starting from the soft polymer with a viscous-like behavior, a significant stress enhancement and a strong reduction of the ultimate elongation are observed with increasing hard polymer loading. As before, this trend is amplified for the coatings dried at 80 ◦ C. This is particularly well illustrated by the Young’s modulus in Fig. 7b, where the reinforcement results in a nearly straight increase for the coatings prepared at 80 ◦ C. The enhancement is less marked for coatings dried at 40 ◦ C, at least at low hard polymer content. Many phenomenological and micromechanical models have been suggested to describe the relation between composition and tensile modulus for immiscible polymer blends and filled polymers [7–11]. This area has been reviewed extensively. A simple upper bound is provided by the parallel or rule-of-mixture (ROM) model predicts a linear relationship for the modulus of the composite or blend material, i.e., Ec,ROM = m Em + f Ef
Fig. 5. TMA indentation profiles of free-standing coatings prepared from the base polyurethanes and their blends. The coatings dried at 40 and 80 ◦ C are shown by solid and dashed lines, respectively. The thickness of the lines indicates increasing contents of hard materials, i.e. 0:100, 20:80, 30:70, 40:60. The hard polyurethane is shown by a double line.
(1)
where m and f = 1 − m stand for the volume fractions of the matrix continuum and the dispersed (filling) compound, respectively, Em and Ef are the tensile moduli of the matrix and the dispersed material, respectively. In this particular case, it is assumed that the two phases deform identically with equal strain. In the opposite limiting situation (lower bound), the two material
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Fig. 6. (a) Composition dependence of the softening temperature. The onset temperature of the major indentation step is shown with large symbols whereas the smaller symbols (diamonds, triangles) are used for the weaker steps. Coatings dried at 40 and 80 ◦ C are represented by solid and open symbols, respectively. (b) High temperature indentation plateau as a function of hard:soft polymer ratio. Coatings dried at 40 and 80 ◦ C are shown by diamonds and triangles, respectively.
components experience equal stress and the composite modulus is then given by Ec,IROM =
Em Ef m Ef + f Em
(2)
This model, known as the inverse rule-of-mixture (IROM), is actually the series equivalent of the parallel model. Fig. 7b shows that the experimental data are contained between the two limits provided by Eqs. (1) and (2). For further comparison, the semiempirical Halpin–Tsai (HT) expression [8] for an isotropic material of spherical particles dispersed in a matrix is also plotted (dashed line). The composite HT modulus reads as Ec,HT = Em
Ef + (m Em + f Ef ) ( + f )Em + m Ef
(3)
where is a parameter related to the geometry and the packing of the filling compound. In Fig. 7b, the dispersed particles were assumed as spherical with a value = 2. The HT equation reduces to Eq. (1) for → ∞ and to Eq. (2) for → 0. Leaving as an adjustable parameter in HT equation does not result in a good fit of the experimental data. The reinforcement effect due to the hard component
grows faster with polymer mass fraction than expected from the Halpsin–Tsai and IROM predictions when volume and mass fractions are considered as equal. This possibly suggests that the hard polymer mass fraction underestimates the effective volume fraction of the dispersed component in the coating. The presence of a compositional gradient (interphase) could partly account for the higher filling volume. However, diffuse phase boundaries are usually beyond the assumptions of the traditional micromechanical models. A notable exception is the self-consistent mechanical model elaborated by Colombini and Maurer which accounts for the occurrence of so-called micromechanical viscoelastic transitions by invoking an interlayer in multiphase polymer materials [25]. Simulations show that an interlayer with sufficient strength can indeed lead to reinforcement but not to the extent that is observed here. Following a simple approach, we have modified the twospring matrix-filler model, represented by Eqs. (1) and (2), into a three compartment model with an additional elastic spring for the interphase region. The experimental data is most conveniently described by a parallel combination of the low (matrix) and medium modulus compartments (interphase) in series with
Fig. 7. (a) Tensile properties of the free-standing films at room temperature. The coatings dried at 40 and 80 ◦ C are shown by solid and dashed lines, respectively. An increasing line thickness refers to an increasing hard:soft ratio. (b) Plot of Young’s Modulus as a function of the hard polymer content in the coating. Drying temperature 40 ◦ C (circles), 80 ◦ C (triangles). The experimental data are compared to the following empirical predictions: ROM (upper dotted line), IROM (lower dotted line), Halpin–Tsai (HT: thick dashed line). The solid lines through the data result from a non-linear least-squares comparison to the three-compartment (TC) model described by Eq. (4).
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Fig. 8. Illustration of some parameters which can be extracted in real time for each tapping cycle by using the PFQNM technique.
the high modulus compartment (filler). The resulting composite modulus reads as f m + int 1 = + Ef Ec Em,int
(4)
where Em,int = (m Em + int Eint )/(m + int ) is the modulus value of the interphase and the matrix in parallel combination. Young’s modulus of the matrix and the filler are, respectively, Em ≈ 0.05 GPa (measured) and Ef ≈ 2 GPa (typical value for a glassy UVPUD). It is assumed that the effective volume fractions f and int are proportional to the mass fraction wf of the dispersed hard polymer, i.e., f = Ff wf , int = Fint wf and m = 1 − f − int . The full lines through the data in Fig. 7b result from a least-squares comparison to Eq. (4) using the proportionality factors Ff and Fint as adjustable parameters. An arithmetic mean value Eint ≈ 1 GPa is tentatively taken for the average modulus of the compartment in-between the filler and matrix compartments. The enhancement of Young’s modulus at low hard polymer content is well captured by this phenomenological model. When comparing the values of the fit parameters at 40 and 80 ◦ C, it appears that the factor Ff reduces from 1.73 to 0.86 whereas Fint inflates from 0.11 to 0.77. These values suggest that the expansion of the interphase at 80 ◦ C is the reason behind the significant enhancement of Young’s modulus compared to the films at 40 ◦ C. The reduction of the filler fraction Ff mainly shifts the divergence to higher mass fractions and has little effect on the behavior at low mass fractions. AFM imaging with Peak Force Quantitative Nanomechanical Property Mapping (PFQNM) [26] is based on the real-time analysis of the force–distance curves recorded at a frequency of about 2 kHz (while usually these curves are recorded at 1 Hz). Since the feedback loop is set on the (peak) force, the actual force is maintained constant during the imaging of the samples. From those force–distance curves, one can extract the following information: (i) the adhesion force (corresponding to the lowest value of the force when the tip is retracted from the sample; (ii) Young’s modulus (obtained from the slope of the curve while the tip is “indenting” the sample, according to the DMT model [27]); (iii) the dissipation (corresponding to the surface integral between the approach and the retract curves); and (iv) the deformation (by comparing the curves to those obtained for a very stiff sample). These parameters are illustrated on Fig. 8. It is important to note that other methods based on a multi-frequency signal sent to the cantilever are also available like Dual AC [28,29], Intermodulation [30] or Band Excitation [13] but it is extremely difficult to extract quantitative values out of the data even if they are providing a contrast in the images. Since the force–distance curves are recorded at every point of the raster scan, the analysis provides a map of those properties. As an example, Fig. 9 illustrates the data obtained for the coating with
a hard:soft ratio of 20:80 at two drying temperatures (left: 40 ◦ C; right: 80 ◦ C). The scale range is 20 nm for the two height images. It is important to note that the topographic images obtained by PFQNM are equivalent to those recorded using TM-AFM. On the adhesion images of the blend with a hard:soft ratio of 20:80, the hard spherical particles are much less adherent than the soft polymer matrix and therefore appear in black in the image. By comparing the roughness of the samples dried at 40 ◦ C (3.25 nm) to the value measured for a drying temperature of 80 ◦ C (1.75 nm), we can conclude that the drying temperature drastically influences the flatness of the dried films. This observation is independent from the hard:soft ratio. More interestingly, the adhesion image for 80 ◦ C (Fig. 9d, right) is less contrasted compared to that of the sample dried at 40 ◦ C and the apparent diameter of the spherical hard particles is smaller. This indicates that the hard spheres are more embedded in the soft film, which is also the main reason why the DMT modulus image is less contrasted. This is fully consistent with the model proposed above, which considers the presence of an interphase between the hard spheres and the soft polymer matrix when the drying temperature is 80 ◦ C. This interpretation is also in agreement with the data from Colombini et al. where they carefully studied the morphology and the viscoelastic properties of bimodal hard/soft latex blends [21]. The latter blends can be considered here as a model compound relative to our UV-PUD systems. 3.4. Film properties of dry UV-cured coatings Regardless of the composition and the drying temperature, all the coatings exhibited a similar thermal behavior after UV-curing characterized by a single glass transition with a Tg value in the range 106–125 ◦ C. The initially soft UV-PUD undergoes a remarkable Tg -increase of more than 100 ◦ C upon UV-curing. Hence, the high crosslink density in both soft and hard regions levels off any compositional contrast, leaving behind an apparently homogenous and extremely hard coating which prevents any indentation experiment. Finally, the tensile behavior is overall hard and brittle with Young’s modulus values of ≈2.2–2.6 GPa and ultimate elongations down to approximately 1%. 3.5. Film performance The soft polymer was compared with the 40:60 blend which was selected as the most promising composition from this study. The polymer dispersions were formulated into orange pigmented compositions for spray application, the choice of the orange color being guided by the fact that it is a challenging color for UVpolymerization owing to the strong competition between the UV-absorption of the photoinitiator and the pigments. For that purpose, they were mixed with an industrial ready-to-use water-based orange pigment paste (15%, w/w) and further formulated with two liquid photoinitiators (1% ESACURETM HB and 0.5% IRGACURETM 819 DW), a rheology modifier (0.5% UCECOATTM 8460) and a wetting agent (1% BYKTM 028). The pigmented formulations were applied using a spray gun technique to MDF panels laminated with white melamine paper and prepared by sanding with a thin aluminum oxide abrasive paper. The coating with a wet thickness of 120–130 g m−2 was dried in a ventilated oven for 20 min at 40 ◦ C and immediately cured on a conveyor belt at a speed of 5 m min−1 using a Ga-doped mediumpressure mercury lamp (power input: 120 W cm−1 ) followed by a medium-pressure mercury lamp (power input: 120 W cm−1 ). The cured coatings were tested for the main performance areas, i.e. adhesion (DIN 53151), stain resistance (EN12720) and nail scratch resistance. The results of the tests are reported on a 0–5 scale (5 = excellent). The adhesion and the stain resistance (black marker, red wine and coffee) of the two colored coatings were all
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Fig. 9. Left: PFQNM images of a thin film of the polyurethane blend with a hard:soft ratio of 20:80. The film was dried at 40 ◦ C. (a) Height; (b) adhesion; (c) DMT modulus; and (d) dissipation. The scan size is 1.0 m. Right: PFQNM images of a thin film of the polyurethane blend with a hard:soft ratio of 20:80. The film was dried at 80 ◦ C. (a) Height; (b) adhesion; (c) DMT modulus; and (d) dissipation. The scan size is 1.0 m. Brighter areas correspond to higher relative values of the image property.
It is clear that the resulting product can be positioned at the high-end performance level for use in pigmented coatings, for instance on wood furniture. It exhibits an excellent colloidal stability and a low viscosity so that it can be easily applied by spray or curtain coating. It is non-irritant and does not contain any cosolvent. It has a low film formation temperature and is tack-free before cure. Importantly, it does not require the use of additional expensive thermal crosslinkers detrimental to the pot-life of the formulation and to occupational health and safety. Acknowledgments
Fig. 10. Comparison of the nail scratch resistance of orange UV-cured coatings based on the soft polymer (0:100) and the blend with the hard:soft ratio of 40:60. The results are plotted as a function of elapsed time (0–5 scale, 5 = best).
found to be excellent (5/5). The results of the nail scratch resistance are shown in Fig. 10 and address the through cure efficiency in the presence of pigments. The performance of the two products increases over the investigated time interval but the blend gives significantly better results immediately after the curing (when the panel is still hot) and reaches a maximum value after 60 min, allowing a good stackability of the coated panels during their manufacture. 4. Conclusions We developed new aliphatic radiation curable polyurethane dispersions by combining soft and hard particles in a blend composition. The blends showed a low film formation temperature associated with a very significant enhancement of the mechanical properties before cure for a hard:soft ratio lower than 50:50. This reinforcement can be explained by the formation of a microstructure with an interphase around the hard spherical particles and is supported by PFQNM AFM data. An optimal ratio was found between the soft and hard components showing at the same time an exceptional chemical and mechanical resistance after curing in the presence of difficult pigments.
The authors wish to thank all the collaborators who were involved with syntheses, application & testing as well as all of the colleagues who gave us their trust and support to develop this new technology. In particular, we are thankful to G. Hellin who made his graduation work both in the research laboratory at CYTEC Surface Specialties and in the laboratory at the University of Mons. We are also grateful to Dr. Chanmin Su (Bruker Nano Inc., Santa Barbara, CA) for helpful discussions on PFQNM. We specially thank Formulaction for their great help in the investigation using ASII. Ph. Leclère is FRS-FNRS Research Associate (Belgium). References [1] K. Buysens, M. Tielemans, Th. Randoux, Pitture e Vernici – European Coatings 19 (2002) 27. [2] M. Tielemans, P. Roose, Ph De Groote, J.-C. Vanovervelt, Prog. Org. Coat. 55 (2006) 128. [3] M. Tielemans, P. Roose, Prog. Org. Coat. 55 (2008) 182. [4] Ph. De Micheli, S. Peeters, Polymer Paint Colour J. 4 (2009). [5] M. Tielemans, J.-P. Bleus, M. Vasconi, Eur. Coat. J. 3 (2007) 38. [6] M. Tielemans, M. Decaux, Ph. De Micheli, S. Peeters, Eur. Coat. J. 4 (2009) 40. [7] R.A. Dickie, J. Appl. Polym. Sci. 17 (1973) 45. [8] J.C. Halpin, J.L. Kardos, Polym. Eng. Sci. 16 (1976) 344. [9] H. Eklind, F.H.J. Maurer, Polymer 37 (1996) 2641. [10] A.G. Facca, M.T. Kortschot, N. Yan, Composites A 37 (2006) 1660. [11] S.-Y. Fu, X.-Q. Feng, B. Lauke, Y.-W. Mai, Composites B 39 (2008) 933. [12] F. Benmouna, T.D. Dimitrova, D. Johannsmann, Langmuir 19 (2003) 10247. [13] S. Jesse, S.V. Kalinin, R. Proksch, A.P. Baddorf, B.J. Rodriguez, Nanotechnology 18 (2007) 435503. [14] R.M. Fitch, Polymer Colloids: A Comprehensive Introduction, Academic Press, London, 1997, pp. 314–337. [15] T. Goverde, U., Boetcher, M. Fleury (Formulaction), private communication.
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