Surface versus Volume Properties on the Nanoscale: Elastomeric

Article pubs.acs.org/Macromolecules

Surface versus Volume Properties on the Nanoscale: Elastomeric Polypropylene Agnieszka Voss, Robert W. Stark,* and Christian Dietz* Department of Materials and Earth Sciences and Center of Smart Interfaces, Physics of Surfaces, Technische Universität Darmstadt, Alarich-Weiss-Str. 10, 64287 Darmstadt, Germany S Supporting Information *

ABSTRACT: The difference between the mechanical properties of a material at the surface and in the bulk is an open issue in polymer science. We studied the mechanical surface properties of polypropylene using atomic force microscopy in peak-force tapping mode. The bulk properties were obtained from layer-by-layer measurements of elasticity, adhesion, and dissipation, with the successive layers removed via wet-chemical ablation. The original sample surface revealed nearly similar mechanical properties for the amorphous and crystalline regions due to a thin (∼22 nm) amorphous top layer. However, in the bulk material, the elastic modulus of crystalline regions was greater than that of amorphous regions. We observed nanoscale crystalline inhomogeneities caused by phase separation that can affect the mechanical stability of polypropylene on the macroscopic scale. The combination of force−volume analysis together with successive ablation of the sample layers form the basis of quantitative nanomechanical tomography.

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INTRODUCTION

The relationship between the macroscopic mechanical behavior and the internal mechanical processes and structure (elastic and inelastic deformation, fracture, etc.) on the microand nanometer scale is still a topic of debate.1,4−7 Systematic design of customized polymeric materials relies on structural and nanomechanical characterization techniques.1,7,8 Most techniques that are used to measure structural and mechanical properties on the nanoscale rely on the atomic force microscope (AFM)9−17 and usually extract mechanical properties, such as the elasticity from force-versus-distance measurements.18,19 Single force-versus-distance measurements, however, provide only point information at a predefined position and hence are not sufficiently dimensioned to understand the correlation between the structural inhomogeneity and the spatial distribution of elasticity or adhesion.20 For full property mapping, various advanced AFM techniques have been proposed to simultaneously acquire the surface topography and maps of the nanomechanical properties of polymeric materials,7,8,20−23 but atomic force microscopy measurements are inherently restricted to the surface. Thus, the volume must be inspected in a layer-by-layer manner to obtain the structural and mechanical information within the bulk of a sample. 4,24−27 This nanotomographic information is of particular interest for semicrystalline polymers because not only the nanomechanical properties but also the number, shape, and spatial arrangement of the crystalline component contribute to the macroscopic elasticity and the material

Recent developments in the synthesis of smart polymeric materials that allow for the modification of various mechanical properties have triggered remarkable scientific and economic interest.1 A prominent example from the broad field of polymers is that of thermoplastic elastomers, also known as thermoplastic rubbers, which are often used in high-impact plastics, pressure-sensitive adhesives, and polymeric foams.2 The functionality of these materials can be customized by finetuning their internal structure. However, to better understand the relationship between the internal structure and the macroscopic material behavior, it is essential to quantify the composition and mechanical properties with nanometer resolution, not only on the surface but also in the bulk material. Polypropylene is a thermoplastic material, and its elastomeric properties depend on the crystallinity of the polymer. Polypropylene with a high degree of crystallinity is an important structural material with outstanding properties, this is, low-cost production, low weight, and high tensile strength, and is widely used in the machinery and automotive industries as well as in electrical and civil engineering. This polymer belongs to the group of polyolefins with a methyl group in each repeating unit. Isotactic polypropylene is able to crystallize because the methyl groups are sterically oriented on the same side with respect to the carbon backbone. In contrast, atactic polypropylene possesses an arbitrary distribution of the methyl side groups along the backbone such that the chains form an amorphous matrix. The elastomeric properties of the material thus can be tuned by balancing the content of atactic and isotactic polypropylene in the semicrystalline polymer.3 © 2014 American Chemical Society

Received: March 20, 2014 Revised: July 22, 2014 Published: July 31, 2014 5236

dx.doi.org/10.1021/ma500578e | Macromolecules 2014, 47, 5236−5245

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strength.1,6,26,28,29 Other central aspects that influence the large scale material properties are inhomogeneities in the crystallinity (i.e., general density fluctuations30), growth of spherulites,31,32 and large-scale differences in the composition and structure caused by liquid−liquid phase separation during the crystallization process.33−35 Additionally, modeling approaches used to predict the viscoelastic behavior of semicrystalline polymers crucially rely on structural information and the nanomechanical properties of the different phases within the material, which are used for model parametrization and validation.29,36,37 In this work, we use the peak-force tapping mode to correlate the morphology of elastomeric polypropylene (ePP) with quantitative mechanical properties, such as elasticity, adhesion, and dissipation. Successive etching provides these physical properties for the layers located beneath the surface. This combination allowed an unprecedented nanomechanical characterization of the volume morphology of elastomeric polypropylene.

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MATERIALS AND METHODS

Peak-Force Tapping. The mechanical surface properties of ePP were quantitatively assessed using the peak-force tapping mode.38 This mode is based on force-versus-distance measurements that are collected with a high acquisition rate of 1−2 kHz (1000−2000 curves/s) to establish two-dimensional arrays containing one forceversus-distance curve at each pixel39 similar to the pulsed force17,40 or jumping mode.41 The cantilever follows the sinusoidal movement of the cantilever base because the system is driven at a frequency far below the resonance frequency of the cantilever. During scanning, the maximum normal force applied to the surface (“peak force”) is controlled by the feedback mechanism of the AFM. Various material properties can be extracted from the retrace portion of the force-versus-distance data. To obtain the elasticity parameters, various models are available based on specific assumptions with respect to the material properties and the active tip−sample interaction mechanisms.42,43 In the peak-force tapping mode, the data (force-versus-distance curves) are fit by the Derjaguin−Muller− Toporov (DMT) model44 to extract the elastic properties of the sample surface (see Figure 1a). This model was developed for weak adhesive forces between the two interacting surfaces and tip radii that are small compared with the compliance of the sample. The loading force FL depends on the deformation δ of the sample surface and can be written as FL =

4 E* R δ 3/2 + Fadh 3

Figure 1. Mechanical surface properties of an unetched ePP sample measured via quantitative nanomechanical mapping: (a) Approach (blue) and retract (red) curve of a single force-versus-separation measurement. The derived physical quantities are highlighted in the graph. (b) Topography image. The bright areas correspond to the crystalline regions of the polymer, whereas the dark areas correspond to the amorphous regions. (c) Corresponding error map for the feedback loop maintaining a constant peak force of 8 nN during imaging. (d) Elasticity map of ePP derived from a DMT model fit. The crystals appear bright due to their high stiffness. (e) Adhesion map of ePP. Please note the inverted contrast. The force necessary to separate the tip from the amorphous regions is greater than the adhesion force on crystalline regions. (f) Map of the energy dissipated between the tip and the sample surface during one oscillation cycle. The determined cantilever/tip properties are k = 6.6 N m−1 and R = 8 nm.

(1)

where E* is the reduced Young’s modulus, R is the tip radius, and Fadh is the adhesion force equal to the pull-off force (i.e., the minimum of the force-versus-distance curve, Figure 1a). Note that δ = z − z0 is the difference between the z-piezo position z measured on the compliant sample and the z-piezo position z0 on an “infinitely” stiff sample at a given load and corresponds only to the deformation of the sample, when z > z0, hence moving the z-piezo further toward the sample surface after contact. In the case of z < z0, δ is the separation/distance between the tip and the sample surface. The reduced elastic modulus E* depends on the elastic moduli of both the tip Etip and the sample Es as well as on their Poisson ratios νtip and νs −1 ⎛1 − ν2 1 − νtip2 ⎞ s ⎟ E* = ⎜⎜ + Etip ⎟⎠ ⎝ Es

polymers are frequency dependent. Also, the adhesion forces can be rate dependent.45 Additionally, the deformation of the sample due to the applied load can be determined from the horizontal difference between the “peak-force” tip−sample separation and the tip−sample separation corresponding to the onset of the overall repulsive force during an approach. The obtained deformation map reveals the local indentation depth of the tip into crystalline and amorphous regions at the maximum applied load. In this study, the DMT model was applied to a heterogeneous polymer exhibiting two physical states, which significantly differ in their mechanical properties. Other contact mechanic models such as the Johnson−Kendall−Roberts46 or Maugis−Dugdale model47 might have been useful to determine more accurate numerical values. It might also have been necessary to switch between models on the different phases. The mathematical complexity of such algorithms, however, does not allow for the implementation into the instrument for an online data analysis at acquisition rates of a few kilohertz. Thus, the absolute numerical values obtained by the fit to the DMT model should be interpreted with the due caution. Imaging and Spectroscopy. Peak-force quantitative nanomechanical measurements (QNM) and single force-versus-distance curves were collected using a Dimension ICON (Bruker AXS, Santa

(2)

The adhesion force is directly accessible from the minimum of the graph in Figure 1a. The hysteresis between the approach and the retrace curve (hatched area of Figure 1a) corresponds to the energy W that dissipates into the sample during one tapping cycle. The elasticity and dissipation values are measured at a specific acquisition rate and do not reflect a full rheological experiment. At higher frequencies the stiffness and dissipation may change as the storage and loss moduli of 5237


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Figure 2. (a) Elasticity map Ek,l, (b) histogram of Ek,l, and (c) determination of crystalline and amorphous elasticity values from the resulting thresholds. (d) Map of the physical state of the polymer 4 = {J , A , C} showing 1 (mixture area) in green, ( (amorphous regions) in blue, and * (crystalline regions) in red. • The amorphous regions are array positions that fulfill

Barbara, CA) atomic force microscope. The NSC-14 (MikroMasch, Wetzlar, Germany) cantilevers with a nominal spring constant k = 5 N m−1 and a nominal tip radius of 8 nm were employed for all experiments. Prior to each measurement, the actual spring constant was determined using the method of Sader et al.,48 and the actual tip radius was estimated by applying the blind tip reconstruction on a sharp-edged TiO2 reference sample (RM15, Bruker AFM probes). Images were collected with a lateral resolution of 512 × 512 pixels. The maximum amplitude was set to 150 nm with a repetition rate of 2 kHz, and the scan speed was set to 1 μm/s. In comparison, single force-versus-distance curves were acquired at a much lower rate of 1 Hz with an amplitude of 450 nm. To extract the elastic modulus from the peak-force tapping data and from the static force-versus-distance curves, we assumed a Poisson ratio of νs = 0.3 for the specimen. The structural and nanomechanical information from multiple layers beneath the surface of the specimen was obtained at the same area. To relocate the area of interest after etching and remounting, the sample was prealigned with the optical microscope by taking advantage of distinctive features on the sample surface. A subsequent enlargement of the scanning area from 10 × 10 μm2 to 3.3 × 3.3 μm2 and finally to 1.1 × 1.1 μm2 as well as permanent realignment aided in finding the same spot repeatedly (compare Figure S2). Data Processing. The images that reflect the mechanical properties of the polymer sample, such as elasticity, adhesion, and dissipation, were calculated by the AFM firmware such that the captured values correspond to the data provided by the AFM. No further image processing was performed on these data. The topography images were first-order flattened to remove sample tilt and to enhance the visibility of surface features. For the roughness analysis of the loading experiment, an offset was removed to preserve the tip-induced structural modification of the surface. The roughness of the polymer surface was measured on a 1 × 1 μm2 image area using the microscope software. Each individual data set Fn, acquired with peak-force tapping after a certain etching step n, can be viewed as a two-dimensional array M × N, which contains a complete force-versus-distance curve Fn,k,l in each element/pixel {(k , l)|k ∈ + : k < M , l ∈ + : l < N }. To distinguish between the crystalline and amorphous regions, we developed an image analysis procedure implemented in MATLAB (MathWorks Inc., Natick, MA) R2011b that analyzes the histogram of the elasticity images. We defined those areas as “amorphous” in which the measured elastic modulus was less than a threshold value, ta, defined as 30% of the maximum relative frequency f max to the left of this maximum in the histogram. Accordingly, the “crystalline” regions were identified by elasticity values above the threshold value tc, defined by the 70% decrease of the maximum relative frequency to the right of this maximum (compare Figure 2). By superimposing the elasticity maps Ek,l for different thresholds with the corresponding elasticity images, the final choices of ta and tc were based on the best match of this overlap. The procedure led to realistic discrimination between the pure amorphous and crystalline regions. Furthermore, each pixel position {(k , l)|k ∈ + : k < M , l ∈ + : l < N } in all of the physical maps (i.e., elasticity map Ek,l, adhesion map Ak,l, and dissipation map Dk,l) was assigned to one of the following classes based on these thresholds: • The crystalline regions are array positions that fulfill

* = {(k , l)|k ∈ + : k < M , l ∈ + : l < N , Ek , l > tc}

( = {(k , l)|k ∈ + : k < M , l ∈ + : l < N , Ek , l < ta}

(4)

• The intermediate class contains all measurements positions that lie in between the amorphous and crystalline regions 1 = {(k , l)|k ∈ + : k < M , l ∈ + : l < N , ta < Ek , l < tc} (5) The result of this classification is a tensor 4 = {1, (, *} that contains the physical states of the polymer, dividing the area of interest into crystalline, amorphous, and intermediate regions. Given a map of any physical property Xn,k,l ∈ {Ek,l, Ak,l, Dk,l} at the etching step n, we can compute the average value of both states X( ̅ and X̅ * as well as their standard deviations ⟨X(⟩ and ⟨X *⟩ with

X( ̅ =

X̅ * =

1 || ( ||

∑ Xk , l

1 || * ||

∑ Xk , l

⟨X(⟩ =

with (k , l) ∈ ( (6)

k ,l

with (k , l) ∈ * (7)

k ,l

1 || ( ||

∑ (Xk , l − X(̅ )2

1 || * ||

∑ (Xk , l − X̅ * )2

with (k , l) ∈ ( (8)

k ,l

and

⟨X *⟩ =

k ,l

with (k , l) ∈ * (9)

In these equations, || ( || and || * || represent the number of entries (pixel) in the corresponding data set. The choice of the threshold parameters tc and ta is a crucial step because the proper classification of the different polymer states depends on it. In Figure 2b, the histogram of the elasticity values of the elasticity image (Figure 2a) is shown. The elasticity values of the stiff crystalline polymer state * and the soft amorphous state ( can be found at both tails of the histogram and the intermediate state 1 that comprises the transition region or mixture of both states is located in the broad center of the distribution. A cutoff value (red line in Figure 2c) was defined at 30% of the histogram’s maximum frequency (percentage of pixel). The respective elasticity on the abscissa are the thresholds ta and tc for the amorphous and crystalline regions. The blue area beneath the curve in Figure 2c (Ek,l < t a ) corresponds to the amorphous regions in the map 4 = {1, (, *} (Figure 2d). Consequently, the red values in Figures 2c,d (Ek,l > tc) correspond to crystalline regions. Please refer to the Supporting Information for complementary information on the data analysis. Sample Preparation. Films of elastomeric polypropylene with a weight-average molecular weight of Mw = 160 kg mol−1 and an [mmmm]-pentade content of 36% (m = meso-conformation) were prepared by drop-casting (100 μL) with a 5 mg/mL ePP/decaline solution on silicon (100). The solution was dried for approximately 24 h in a fume hood at a temperature of 24 ± 2 °C, resulting in ePP films with thicknesses of >1 μm. Etching. Thin layers of ePP were successively ablated by wetchemical etching with potassium permanganate dissolved in sulfuric

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Figure 3. Etching series depicting the mechanical properties of ePP in the volume. The left column shows the mechanical properties of ePP after one etching step, the central column corresponds to the ninth etching step, and the right column presents the summarized mechanical properties of each layer during etching: (a) topographical image and roughness of the sample; (b) elastic modulus maps and elasticity trend during etching for crystalline (red diamonds) and amorphous (blue circles) regions derived from the DMT model. The respective maps and trend of adhesion (c) and dissipation (d) versus etching step for crystalline (red diamonds) and amorphous (blue circles) regions are shown. Each etching step caused removal of an approximately 11 nm thick surface layer. acid according to ref 26. An amount of 1 g of potassium permanganate was mixed in 20 mL of a 30 wt % sulfuric acid for 10 min at a stirring rate of 300 rpm. The sample was dipped into the etching solution for 1 min and subsequently rinsed with a 10 wt % sulfuric acid, hydrogen peroxide, pure water, and acetone, followed by drying in a nitrogen flow; this process caused a removal of approximately 11 nm of the top layer. The etching rate was derived by measuring the film thickness of a reference sample before and after subsequent etching.26

crystalline (bright) and amorphous portions (dark) of the semicrystalline polymer. The root-mean-square (rms) roughness of the surface was approximately 4.0 nm. Figure 1c shows the corresponding error map (peak-force error) for the feedback loop maintaining a constant peak force of 8 nN during imaging. The elasticity map (Figure 1d) reveals the elastic modulus of the crystalline and amorphous regions as 154 ± 10 and 91 ± 2 MPa, respectively. Local maps of adhesion and energy dissipated between the tip and the polymer sample are visualized in Figures 1e and 1f, respectively. The adhesion values for the crystals were 9.6 ± 0.9 nN, and the adhesion in the amorphous regions was 12.8 ± 0.4 nN. The respective values for the dissipation were 0.8 ± 0.2 keV on the crystals and 2.3 ± 0.3 keV on the amorphous regions. Note that the contrast

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RESULTS AND DISCUSSION The diversity of the mechanical properties of elastomeric polypropylene, which were derived pixel-wise from the forceversus-distance data, as illustrated in Figure 1a, is illustrated in Figure 1b−f. In the topographical image (Figure 1b), two distinctive regions are visible, which can be allocated to the 5239


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top layers of the specimen (Figure 3b,c). However, the dissipation maps show similar features for the full etching series (Figure 3d). This observation indicates a good stability of the tip and a stable lateral resolution for the images shown. The chemical treatment during the ablation process potentially affects the measured volume properties. After etching and removing an 11 nm thick layer there might be residual chemical modifications affecting the topmost remaining molecules. The indentation depth during peak-force tapping, however, was a few nanometers. Thus, we assume that the main contribution to the elasticity signal came from the compressed volume beneath the tip and only a small contribution from the potentially modified surface. To quantify the physical properties of the sample surface, we distinguished between areas containing crystalline and amorphous regions, as described in the Materials and Methods section and in the Supporting Information. Intermediate regions were not considered for this analysis. The values representing the physical properties of both areas are plotted for each etching step in the graphs of Figure 3 (right column). The rms roughness of the surface (Figure 3a, right) increases with ongoing etching and saturates after the fourth etching step at approximately 12 nm. After the sample preparation, the DMT model fit yielded an elastic modulus of 91 ± 2 MPa on the amorphous and 154 ± 10 MPa on the crystalline regions (Figure 3b, right). Removing the first two layers (∼22 nm) led to an increased DMT modulus on the crystals, which saturated at a value of approximately 548 MPa. In contrast, the DMT modulus found on the amorphous regions increased only slightly and saturated at approximately 256 MPa. The adhesion force between the tip and the polymer was nearly constant for both regions (Fadh,cr = 6.5 nN, Fadh,am = 12.3 nN) within the variation of the standard deviation for etching through the volume of the sample except for the top layer. In this layer, we obtained an adhesion force on the crystalline regions that was within the range of the standard deviation of the adhesion found on the amorphous regions. We observed a continuous decrease from 2.3 to 1.4 keV of the energy dissipated between the tip and the sample surface (per approach-and-retract cycle) on the amorphous regions. On the crystalline regions, however, the dissipation is smaller (0.8 keV) and remains constant after an initial drop with ongoing etching (0.1 keV on average). Note that all measurements were performed with a peak force of 8 nN. The increasing rms roughness up to the fourth etching step is a consequence of differences in the etching rate between the amorphous and the crystallized polymer.26 For further etching, the differences are balanced by the increased amount of crystalline portions within the area of interest. The nearly identical values of the elasticity and adhesion for both regions indicate the presence of a thin amorphous layer on top of the crystalline regions for an as-prepared sample, and these values corroborate the findings of earlier studies.25,51 This layer was removed after two etching steps, thus exposing the crystal surface. The elasticity of the amorphous phase agrees well with the values reported by Gracias et al.50 Nevertheless, the measured stiffness of the crystals was smaller. The discrepancy can be explained by the semicrystalline nature of elastomeric polypropylene in which crystals are embedded in an amorphous matrix, thus lowering the effective stiffness. The viscoelastic nature of the amorphous polymer becomes dominant in the high dissipation values per oscillation cycle compared with that of the crystallized polymer. The adhesion values on the

in these images is inverted compared with that of Figure 1b,d due to the higher adhesion and dissipation values measured on the amorphous portions compared with the crystalline portions of the polymer. Peak-force tapping also allows for the measurement of the sample deformation. In contrast to the elasticity, adhesion, and dissipation values, however, it is very difficult to obtain reliable indentation values with an automated procedure in this case. The morphology exhibits similarities to the bundled flowerlike structure at a later stage of crystallization, as described by Schönherr et al.49 Our specimen showed a higher density of crystals at the surface than that of Schönherr et al. due to the higher content of [mmmm]-pentade of the polymer used in this work (36%). In addition, our samples were measured 1 week after preparation, when most of the isotactic polymer chains are expected to be in the crystal phase. The surface roughness of 4 nm is a consequence of the difference in the sample deformation/indentation due to the applied load of the tip between the crystalline and amorphous regions. This suggests that the sample surface is rather flat, and the measured roughness is mainly tip-induced instead. The localized DMT modulus measured on the amorphous portions of ePP was on the same order of magnitude as the values reported by Gracias et al.50 using tips with large radii of curvature (1 μm). The surprisingly low DMT modulus measured on the crystalline regions, however, can be explained by the presence of a thin amorphous layer covering the surface of the crystals.25,51 Considering the geometric tip−sample convolution and the widths of single crystalline lamellae of 14 nm in the αmodification and 7 nm in the γ-modification,52 the lateral resolution of the local elasticity map is striking because extremely thin crystalline connections (∼15 nm; see red arrow in Figure 1d) are visible in the DMT modulus image. The high adhesion force and dissipated energy measured on the amorphous phase compared with those of the crystalline portions can be ascribed to the viscous properties of an amorphous polymer in the nonglassy state.53 The visibility of the crystals in the topography image before etching can be attributed to the feedback loop mechanism that solely triggers on the peak-force set-point. The indentation of the tip into the sample surface enables to sense the crystals even when they are buried beneath a thin amorphous top layer. To obtain the mechanical properties of ePP beneath the surface, we successively removed thin layers of the surface and gathered the accessible information on the surface roughness (Figure 3a), elasticity (Figure 3b), adhesion (Figure 3c), and dissipation (Figure 3d) after each etching step. Figure 3 explicitly shows maps of these sample properties after the first etching (left column) and after nine etching steps (middle column), corresponding to sample depths of approximately 11 and 99 nm, respectively. The entire set of images of the topography and elastic modulus values for the complete etching series is shown in Figures S3 and S4 of the Supporting Information. After the first etching, the bundled flowerlike crystalline structure, which was predominant at the sample surface, was ablated (Figure 3a, left). Thereby, the percentage of the crystalline area with respect to the entire surface area was increased. This effect is apparent in all of the maps accessible by the peak-force tapping mode (Figure 3a−d, left column). After the ninth etching step, the crystalline regions were nearly indistinguishable from the amorphous areas in the topographical image (Figure 3a, middle), but the microstructure of the crystalline regions is similar fine-structured as that of the 5240


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min−1. Note that the corresponding elasticity image gives an optical impression of the high crystallinity. In general, the relationship between the two observables demonstrate that the greater the crystallinity within the defect area, the greater the amount of material removed in the subsequent etching step of the same area. Considering the complete etching series (see Supporting Information, Figures S3 and S4), we estimate a depth of 30−40 nm for the defect within the volume structure of the specimen because there was no significant change either in the crystallinity or in the depth of the hollow after the fourth etching step. We postulate that the origin of these defects lies in both the polydispersity of the stock solution and the sample preparation. The polydispersity of the polymer implicates a phase separation between short and long polymer chains in the liquid solution.34 Different types of polymers show a similar behavior.33,54 From a thermodynamic point of view, short isotactic chains are able to crystallize faster than long isotactic chains at low temperatures.55 In contrast with the sample preparation from the melt, the quick evaporation of the solvent additionally accelerates the crystallization process. As a consequence, the probability of embedding noncrystallizable atactic chains into the crystal lattice during the crystallization process by means of tensile strain is high and can lead to a variety of lattice defects.56,57 These lattice defects in turn, i.e., atactic chains, offer a point of attack for the etching solution. The crystal might crumble during etching, which explains the high etching rate and hence the development of hollows on the sample surface despite the high content of crystals within these defect areas. It is of major interest to understand how the polymer is affected by the applied load of the tip during imaging to verify the values obtained by peak-force tapping. Hence, we tested the robustness of the polymer surface by successively scanning a specific area, increased the applied load from Fp = 4 nN to Fp = 120 nN, and, subsequently, compared the physical properties with the properties of the vicinity after each image. To remove the amorphous top layer, we etched the sample twice. Figure 5a compares the rms roughness for different applied loads of the influenced (1 × 1) μm2 area (upward pointing triangles enclosed by the white dashed frame in Figure 5b) with the rms roughness of a reference area (downward pointing triangles) of the same size, as chosen in the lower left corner of the image. Within the applied range of peak forces, increasing the force lead to an increased roughness of the influenced area. A kink is clearly visible at a peak force of Fp ≈ 44 nN, where the graph changes slope. In contrast, the roughness of the reference area remained constant during imaging of the overview scans with a peak force of Fp = 3 nN. Figure 5c depicts the energy dissipated between the tip and the different polymer regions during one approach-and-retract cycle. At small peak forces, the energy per cycle dissipated between the tip and the crystalline polypropylene (diamond symbols) increases only slightly with increasing load. At a peak force of Fp = 44 nN, this increase was stronger. The energy per cycle dissipated between the tip and the amorphous polypropylene (round symbols) increases with a constant slope comparable to the slope measured on the crystalline regions for Fp > 44 nN. A permanent imprint caused by the tip becomes apparent in the series of overview scans of Figure 5b. The respective averaged cross-sectional profiles (3 × 1 μm2) are shown in Figure 6. The black line depicts the original profile. Increasing the force to Fp = 44 nN caused a small imprint within the affected area. The two dashed lines illustrate the boundaries of the affected area. However, a further

crystalline regions in the volume of the sample were rather high for the interaction between a pure SiO2-tip and a polypropylene surface in the crystal state. There might be an adsorption/ attachment of amorphous polymer chains during imaging which increases the obtained adhesion values. The measured elasticity, however, remains unaffected by this very thin polymer layer covering the tip. During the layer-by-layer ablation of the sample surface, we observed small regions that were etched more rapidly than the surrounding area. In the following, we denote these areas as defect areas/regions. These spherical areas were distributed over the entire sample but represented only a minor fraction of the complete surface (see Figure S2). To further investigate this phenomenon, we focused on a single defect region and established the dependency of the relative etching rate on the crystallinity of this area with respect to its surroundings (Figure 4). We defined the relative etching rate as Δr = rd − rs =

href − hd h − hs h − hd − ref = s Δt Δt Δt

(10)

Figure 4. Dependence of the relative etching rate of a defect area (open red diamonds) on the crystallinity with respect to the surrounding area. The insets show the respective height and elasticity images for etching steps 1−4 (red numbers note the order). The dashed white frames enclose the area of crystals with structural defects. High crystallinity within the marked area leads to a high etching rate compared with the surrounding area in the following etching step. As a consequence, a hollow arises in this region.

where rd and rs are respectively the etching rates of the defect and surrounding area, hd and hs are the actual respective surface levels after the etching of both areas with respect to the initial sample thickness href as a reference level, and Δt is the time we set for one etching step. The crystallinity of the defect area is the number of pixels belonging to crystals (see Data Processing section) divided by the total number of pixels within the defect area (dashed white frame in Figure 4). After the first etching step (the data points for the etching steps are designated by Arabic numbers in the graph), we measured a crystallinity of 6.7% and a relative etching rate of 2 nm min−1 for the forthcoming etching step; i.e., on average, 2 nm was additionally ablated from the defect region compared with the rest of the sample surface. The inset connected to each data point shows the corresponding topography image and the map of elasticity. After etching step number 2, the crystallinity was determined as 11.1%, leading to a relative etching rate of 15 nm 5241


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Figure 6. Averaged cross-sectional profiles (averaged area of 3 × 1 μm2). The profiles were obtained after surface modification caused by an increased surface load. The gray lines indicate the boundary of the region of interest (area of increased loading force), as indicated in Figure 5. The black line corresponds to the surface profile measured after an applied load of Fp = 4 nN, the gray line to Fp = 44 nN, and the red line to Fp = 120 nN. The primary scan direction was from left to right.

Figure 5. Influence of the tip on an etched ePP surface: (a) Progress of the rms roughness with increasing load between the tip and the ePP sample. Open black triangles (pointing up) correspond to the rms roughness values of the 1 × 1 μm2 area marked with a white dashed frame shown in the topographical images (b). The gray open triangles (pointing down) correspond to a 1 × 1 μm2 reference area of the lower left part of the same images. Plastic deformation during this imaging series of increasing load becomes clearly visible in the topography. Arabic numbers link the images to the roughness diagram. The increased forces were only applied to the framed region of interest. The images were measured with a peak force of 3 nN only. Polymeric material was removed from the left to the right direction. Please refer to Figure S5 of the Supporting Information for the corresponding elasticity and adhesion maps of the same area. (c) Energy dissipated between the tip and the sample during one approach−retract cycle measured on the crystalline (red open diamonds) and the amorphous regions (blue open circles) within the marked area. Linear regressions were fit to the roughness and dissipation values corresponding to the peak-force values of Fp < 44 nN and Fp ≥ 44 nN. The gray dashed lines indicate the intersection of the fitted curves at which plastic deformation of the polymer occurred for the first time. The determined cantilever/tip properties are k = 8.6 N m−1 and R = 2 nm.

effect occurred at high loads that enhanced the surface roughness. As shown in the cross-sectional profiles of Figure 6, polymeric material was moved in the direction of the primary scan. Most likely, this material predominantly consisted of loose atactic polymer chains not embedded into the crystal structure of isotactic chains and hence sufficiently mobile for relocation. The two explanatory mechanisms (pushing crystals, moving material) are opposing trends with a stronger impact of the material movement on the surface roughness. This observation explains the reduced slope of the surface roughness with increasing peak-force for Fp > 44 nN vs Fp < 44 nN. The increase of the energy dissipated between the tip and the amorphous polymer (Figure 5c) with increasing force is typical for viscoelastic properties. For peak-forces below Fp < 44 nN, the energy dissipated between tip and crystalline regions remained almost constant. To a certain degree, peak-force tapping can be compared with dynamic atomic force microscopy modes in which similar behavior was found using spectroscopy techniques.59 The increase in energy dissipation with increasing peak-force for Fp > 44 nN corroborates our postulation of an existing soft amorphous matrix in which single crystalline lamellae can be pushed because the slope of the dissipation versus peak-force exactly matches the slope found on the amorphous regions. We emphasize that the applied pressure is the key parameter that determines whether the measurement is nondestructive/ destructive due to shear stresses in the amorphous polymer. Calculating the exact pressure from the applied force and the contact area, however, is challenging because the contact radius depends on the indentation and hence on the applied force. Measuring the tip radius before and after the experiment and comparing the resolution of the surrounding area of the overview scans indicate that the change in sample surface within the influenced area visible in Figure 5b is caused by the destruction of the sample surface rather than by tip wear. It is well-known that the viscoelastic properties of many polymers depend on the frequency/shear rate in rheological experiments. We compared peak-force tapping conducted at a relatively high acquisition rate (2 kHz) with conventional

increase of the applied force to Fp = 120 nN caused a hollow on the left side of the imaged area and an elevated feature on the right side to appear (highlighted by arrows in Figure 6). The trajectory of the tip in the direction of the fast scan axis was from left to right (primary scan direction) and back (secondary scan direction). The experiment implies a threshold that separates the elastic from the inelastic behavior of the polymer on the nanoscale. In the elastic range, the more pressure we applied to the polymer surface, the greater the indentation of the tip into the amorphous polypropylene. In the crystalline regions, however, the indentation/deformation remained constant with increasing force. As a consequence, the crystals increasingly protruded from the amorphous surroundings, leading to a higher apparent surface roughness.27,58 As it exceeded the threshold force (Fp > 44 nN), the tip began to push the crystals irreversibly into the amorphous vicinity, which served as a soft matrix. One could argue that embedding the crystals into the soft matrix would consequently decrease the surface roughness. A secondary 5242


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“quasi-static” force-versus-distance measurements (accomplished at 1−2 Hz). The fast penetration of the tip into the sample surface during a load−unload cycle for peak-force tapping might be considered as a dynamic shear experiment. To assess the obtained data, we randomly compared the elasticity values obtained by peak-force tapping with values obtained by applying the DMT model (eq 1) to the “quasi-static” forceversus-distance measurements. Furthermore, we manually fitted the DMT model to single force-versus-distance curves from the peak-force-tapping data set and extracted the elastic modulus. Prior to the measurement, the amorphous top layer was removed by wet-chemical etching. The DMT moduli provided by the operational software of the AFM and the subsequent data analysis (see Data Processing section and Supporting Information) for the image shown in Figure 7a resulted in an average value of approximately Eam = 174 MPa on the amorphous regions and Ecr = 489 MPa on the crystalline regions. The red diamond-shaped data points and the blue circles respectively represent the retrace regions of the forceversus-distance curves measured on crystalline and amorphous regions extracted from the peak-force tapping data set at the positions marked with the crosses. We fit the DMT model to these curves as indicated by the black lines and obtained elastic moduli of Eam = 133 MPa and Ecr = 271 MPa for amorphous and crystalline polypropylene, respectively. Single force-versusdistance measurements (Figure 5b) at a low acquisition rate (1 Hz) and the associated DMT fit resulted in Eam = 49 MPa and Ecr = 312 MPa. The compared methods resulted in only slightly different elasticity values. To understand the mechanism of the tip− sample interaction of each type of measurement, we considered the trajectory of the tip and the velocities with which the tip touches the surface (immersion speed). The black line in Figure 7c shows the progress with time of the tip position in the z-direction for a sinusoidal excitation as applied in the peakforce tapping mode. In this figure, the sample surface is illustrated as a horizontal gray line for an indentation depth of 6 nm, which is typical for amorphous regions for the experimental conditions applied in this work. The intersections of both lines are the points in time at which the tip is immersed into the polymer and subsequently detaches from the surface after one period of interaction. From the first intersection, we determine an immersion speed vpf = 0.5 mm/s by calculating the derivative of the sinusoidal function (red line) and extracting the respective velocity (indicated as red open dots in Figure 7c). For static force-versus-distance measurements, we consider a temporal triangular function performed by the tip. This leads to a constant velocity of vstat = 1 μm/s for the chosen experimental parameters (dashed red line). Consideration of the notably large difference in the immersion velocities (greater than 2 orders of magnitude) of the forceversus-distance measurements acquired at 1 Hz and 2 kHz and the good agreement of the obtained elasticity values leads to the conclusion that only a weak dependency exists in the elasticity of both components (crystalline and amorphous) in the investigated range of shear rates. Elastomeric polypropylene is a composite material consisting of an amorphous and a crystalline material with different mechanical properties. However, the mechanical properties of each component within the composite differ from that of pure isotactic or atactic polypropylene.50,60

Figure 7. Comparison of force-versus-separation (retrace) curves extracted from (a) a peak-force-tapping image (acquisition rate of 2 kHz) with (b) manually conducted “quasi-static” force-versusseparation curves (1 Hz). The red open diamond symbols correspond to force values measured in crystalline regions, whereas the blue open circle values were measured in the amorphous regions. The crosses in the insets (elasticity maps) mark the points where the curves were taken. The black lines represent the DMT fits for each curve within the 30−90% region related to the total force range applied to the sample. The determined cantilever/tip properties are k = 7.7 N m−1 and R = 9 nm. (c) Illustration of the tip movement (black line) and the velocity (red lines) as a function of time for the peak-force-tapping mode. The intersections of the black line with the sample surface illustrated in gray lines represent the points at which the tip first touches the surface and loses contact after indenting. These points can be projected (gray dashed lines) onto the velocity curve (red line) to reveal the immersion speed of the tip into the material. The red dashed line represents the constant velocity during an approach-and-retract cycle of the quasi-static force-versus-distance curves. Arrows link the curves to the respective y-axis.

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CONCLUSION We investigated the mechanical properties of elastomeric polypropylene using an atomic force microscopy technique 5243


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Universität Chemnitz for fruitful discussions and for sample courtesy. We also thank the Center of Smart Interface for financial support.

based on force-volume imaging that allowed for the simultaneous mapping of the elasticity, adhesion, and dissipation, with a lateral resolution of approximately 15 nm. The elastic modulus, adhesion, and dissipation values measured on crystalline regions were close to the values on the amorphous regions for untreated samples. Successive ablation of the surface led to different values for these mechanical material properties, implying that a thin amorphous polymer layer initially covers the crystalline regions. In general, the amorphous polypropylene was removed faster than the crystalline polypropylene in wet-chemical etching except for regions characterized by a conspicuously high crystallinity, which appeared as defect regions. This inhomogeneity caused hollows on the sample surface during ablation. We postulate that these defect regions primarily consist of short isotactic polymer chains that phase-separated from the long chains in the liquid phase and crystallized rapidly but impurely. Applying a maximum force of more than 44 nN during imaging caused plastic deformations to the sample surface. Imaging the polymer successively with an increasing force resulted in a material dislocation along the primary scan direction of the tip. Consequently, a hollow on one side and an elevated feature on the other side occurred within the imaged area. Reasonably, the transported material primarily consisted of loose unbound atactic chains of polypropylene. A comparison of the immersion speeds of the tip into the polymer for force-versus-distance measurements accomplished at 2 kHz (peak-force tapping mode) and 1 Hz (conventional force-versus-distance measurements) was drawn. Because of the difference of more than 2 orders of magnitude in the velocity upon the contact of the tip with the surface and the similarity of the obtained values for the elasticity for both techniques, we conclude that only a weak dependency exists for amorphous and crystalline phases of polypropylene in the elastomeric configuration on the acquisition rate of the conducted experiment. In summary, the combination of peak-force tapping and layer-by-layer etching of elastomeric polypropylene permitted us to find a thin amorphous layer on top of a freshly prepared film as well as regions of structural defects in the crystalline component of the material, which also appeared in the volume of these films. This inhomogeneity measured on the nanoscale can affect the mechanical stability of polypropylene on the macroscopic scale. The technique presented in this work can be applied to other classes of material and provides a basis for full quantitative nanomechanical tomography when combined with nanotomography.24

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ASSOCIATED CONTENT

S Supporting Information *

Data processing; Figures S1−S5. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (R.W.S.). *E-mail [email protected] (C.D.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge Eike-Christian Spitzner, Martin Neumann, and Robert Magerle from the Technische 5244


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