Evaluating the Reinforcement of Inorganic Fullerene-like

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Evaluating the Reinforcement of Inorganic Fullerene-like Nanoparticles in Thermoplastic Matrices by Depth-sensing Indentation Araceli Flores, Mohammed Naffakh, Ana Maria DíezPascual, Fernando Ania, and Marián Angeles Gómez-Fatou J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/jp406513y • Publication Date (Web): 17 Sep 2013 Downloaded from http://pubs.acs.org on September 20, 2013

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The Journal of Physical Chemistry

Evaluating the Reinforcement of Inorganic Fullerene-like Nanoparticles in Thermoplastic Matrices by Depth-sensing Indentation Araceli Flores,†* Mohammed Naffakh,‡ Ana M. Díez-Pascual,§ Fernando Ania,† Marián A. Gómez-Fatou§ †

Departamento de Física Macromolecular, Instituto de Estructura de la Materia (IEM-

CSIC), Serrano 119, 28006 Madrid, Spain ‡

Universidad Politécnica de Madrid, Departamento de Ingeniería y Ciencia de los

Materiales, Escuela Técnica Superior de Ingenieros Industriales, José Gutiérrez Abascal 2, 28006 Madrid, Spain §

Departamento de Física de Polímeros, Elastómeros y Aplicaciones Energéticas, Instituto

de Ciencia y Tecnología de Polímeros (ICTP-CSIC), Juan de la Cierva 3, 28006 Madrid, Spain *Corresponding author: [email protected]. Phone: + 34 917459523

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ACS Paragon Plus Environment


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ABSTRACT: The reinforcing effect of inorganic fullerene-like tungsten disulphide (IFWS2) nanoparticles in two different polymer matrices, isotactic polypropylene (iPP) and polyphenylene sulphide (PPS), has been investigated by means of dynamic depthsensing indentation. The hardness and elastic modulus enhancement upon filler addition is analyzed in terms of two main contributions: changes in the polymer matrix nanostructure and intrinsic properties of the filler including matrix-particle load transfer. It is found that the latter mainly determines the overall mechanical improvement, whereas the nanostructural changes induced in the polymer matrix only contribute to a minor extent. Important differences are suggested between the mechanisms of deformation in the two nanocomposites resulting in a moderate mechanical enhancement in case of iPP (20% for a filler loading of 1%), and a remarkable hardness increase in case of PPS (60% for the same filler content). The nature of the polymer amorphous phase, whether in the glassy or rubbery state, seems to play here an important role. Finally, nanoindentation and dynamic mechanical analysis measurements are compared and discussed in terms of the different directionality of the stresses applied.

Keywords: Instrumented Indentation, Polymer Nanocomposites, Elastic Modulus, Hardness, WS2 Nanoparticles, Continuous Stiffness Measurements.

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1. INTRODUCTION Inorganic fullerene-like (IF) nanoparticles and inorganic nanotubes (INT) based on layered metal dichalcogenides, such as WS2 and MoS2, have emerged as one of the most promising developments in the area of nanomaterials. The first synthesis of such nanostructures was reported by Tenne et al. in 1992 and 1993.1,2 Ever since, the number of reports on the successful growth of IFs from inorganic compounds has increased rapidly, emphasizing the importance of this field for nanotechnology.3,4 Concurrently, potential applications in catalysis, rechargeable batteries, drug delivery, solar cells and electronics have been proposed.3 In addition, these kind of nanostructures represent promising candidates as part of ultrahigh-strength polymer nanocomposites5 due to its superior mechanical properties like very high stiffness and strength.6 Inorganic fullerene-like nanoparticles such as those made of tungsten disulphide (IF-WS2) exhibit particle sizes typically in the range of 40-180 nm, quasispherical shape, closed-cage layered structure and chemical inertness. They are cheaper than organic nanofillers (i.e. carbon nanotubes, nanofibers, graphene), more environmentally friendly and display a lubricant characteristic that facilitates their homogenous dispersion within polymer matrices up to high loadings.5 By combining the merits of organic and inorganic (IF) materials in novel polymer/IFs, different types of nanocomposites and hybrid materials have been successfully developed ranging from high-performance to commodity thermoplastic-based systems, making them an exciting research area for the next years.5,7-12 IFs can also be used with other organic micro-particles (nucleating agents), micro-fibers (carbon fibers) or nanofillers (carbon nanotubes) as innovative strategy to obtain advanced hybrid materials with complex architectures, interactions, morphologies and functionalities.5 IF nanoparticles have been shown to be efficient for improving the thermal, mechanical and tribological properties in a number of thermoplastic polymers, including isotactic polypropylene (iPP),7 polyphenylene sulfide (PPS),8-10 poly(ether ether ketone) 3



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(PEEK),11 and nylon-6,12 that were processed via traditional melt processing techniques as economical and scalable routes, without the need of modifiers, surfactants or dispersing agents. In particular, in the case of iPP based nanocomposites,7 the nanoparticles readily acted as nucleating agents increasing the crystallization rate of the matrix, their nucleating efficiency reaching very high values (60–70%), the highest reported for iPP nanocomposites. In contrast, a drastic change from a delay to a boost in the crystallization rate was observed for PPS based nanocomposites when the nanofiller content increased from 0.1 to 0.5 wt%.9,10 On the other hand, the nanoparticles induced remarkable enhancements in the thermal stability and mechanical performance of the composites, showing increases in the storage modulus in the range of 7-25% and 40-75% for iPP7 and PPS10 based composites, respectively. Instrumented indentation (or depth-sensing indentation, DSI) has gained an outstanding relevance in the last decade as a means of characterizing the mechanical properties of materials. The attractiveness of the technique is based on the ability to probe small volumes of materials (the size of the contact depth could be as low as only a few nanometers). Analysis of DSI data using well-established methods provides hardness, H, and elastic modulus, E, values.

13,14

However, important limitations arise in case of polymer

materials due to their time-dependent character.15-18 The advent of the continuous stiffness method (CSM) that introduces a small oscillating force during the loading cycle, represents a step forward in the study of viscoelastic materials in general, and polymers in particular, by means of DSI.

14

Accordingly, modulus and hardness can be achieved through the whole

loading cycle, i.e., as a function of indentation depth. Whether or not the modulus values obtained by dynamic indentation testing agree with those obtained by conventional bulk techniques such as Dynamic Mechanical Analysis (DMA), uniaxial compression or tensile testing is still a matter of debate. 19-25

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Meanwhile, instrumented indentation has been successfully applied to the study of the effect of reinforcement in polymer nanocomposites containing different fillers such as nanoclays, carbon nanotubes, graphene, nanodiamond, etc.23-29 This technique has been proved to be sensitive to filler concentration and dispersion and, in addition, to the interfacial nanofiller-matrix adhesion.23,27 Morphological changes taking place in the polymer matrix as a consequence of the incorporation of the filler24,27 and the homogeneity of the composite specimen across the thickness24,25 are also topics of research by means of DSI. The aim of the present study is to evaluate the reinforcing effect of inorganic fullerene-like particles in two different polymer matrices by means of DSI and to compare the results with those obtained by DMA. The nanofiller quasi-spherical shape presents two main advantages: on the one hand, it does not induce an average preferential orientation to the neighboring polymer chains; on the other hand, spherical nanoparticles should not preferentially orient along the direction of the applied load in uniaxial tensile testing, hence, facilitating the comparison between DSI and tensile DMA studies. The two polymer matrices were selected to exhibit their glass transition temperatures well separated from each other, one above and one below the test temperature (room temperature), in order to explore the influence of the nature of the amorphous phase on the filler reinforcing effect. In addition, because both PP and PPS are semicrystalline matrices, an analysis of the mechanical enhancement upon addition of the nanofiller is offered with the aim of separating the contribution of the filler reinforcement itself and that associated to changes in the matrix nanostructure induced by the presence of the filler.

2. EXPERIMENTAL SECTION 2.1. Materials iPP was supplied by Repsol-YPF (Spain), with 95% isotacticity, a viscosity average molecular weight of 179,000 g/mol and a polydispersity of 4.77. PPS 5



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(Fortron 02054P4) was supplied in pellet form by Ticona (Spain). Inorganic fullerene-like tungsten disulfide (IF-WS2) nanoparticles (NanoLubTM, d25ºC ~ 7.5 g/cm3, mean diameter of 80 nm) were provided by Nanomaterials (Israel). Different concentrations of IF-WS2 (0.058.0 wt%) were introduced in the thermoplastic matrices by melt-blending in a Haake Rheocord 90 extruder operating at 210 and 320 ºC for iPP and PPS, respectively, with mixing times of 20 min and a rotor speed of 150 rpm. The samples were subsequently compression moulded to obtain films ~ 0.2 mm thick using a hot-press system (Collin) with two pairs of heating/cooling plates (the melt temperature was 210 ºC for iPP and 320 ºC for PPS, time = 5 min and pressure = 13 MPa). In particular, the amorphous PPS film was prepared by hotcompression at 320 ºC followed by immediate quenching in liquid nitrogen. 2.2. X-ray Diffraction Measurements Room temperature wide angle X-ray diffractograms were obtained in the angular range of 2θ = 5º – 40º with a Bruker D8 Advance diffractometer using a Cu tube as X-ray source (λ CuKα = 1.54 Å), with a voltage of 40 kV and an intensity of 40 mA. The degree of crystallinity ωc of the samples was calculated using the following relation: ωc = Ic / (Ic + Ia), being Ic and Ia the integrated intensities of the crystalline and amorphous phases, respectively. The average lateral crystal size of the composites in the (110) direction (D110) was calculated from the diffraction patterns following the Scherrer equation: Dhkl = λ/β cosθ, where β is the full width at half maximum (in rad) of the crystalline peak.30 2.3. Dynamic Mechanical Analysis The

dynamic-mechanical

performance

of

polymer nanocomposites was studied using a dynamic mechanical analyzer (Mettler DMA861). The samples were cut in rectangular shapes (PP ∼ 19.5×4×0.5 mm3, PPS ∼ 19.5×4×1.0 mm3) and mounted in a large tension clamp. The measurements were carried out in the tensile mode, in a temperature range (PP=-100 to 100 ºC, PPS=-100 to 230 ºC) and with

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a heating rate of 2 ºC/min. A dynamic force of 6 N oscillating at a fixed frequency of 10 Hz and 30 µm of amplitude was used. 2.4. Nanoidentation Tests

Portions of the nanocomposite films were glued onto a

metallic holder that was placed in the stage of the Nanoindenter. DSI tests were performed on a Nanoindenter G200 (Agilent Tech.) using the continuous stiffness measurement (CSM) technique. Experiments were carried out in depth-control mode, using a maximum indentation depth of 5 µm. During the loading cycle, the load was incremented at a constant P´/P ratio in order to ensure a constant indentation strain rate (0.2 s-1) during the loading cycle. A sinusoidal force of 2 nm of amplitude and 45 Hz of frequency was superimposed to the quasistatic load during the loading. The indenter displacement is phase-shifted with respect to the excitation force. Using a simple harmonic oscillator analogy to model the combined dynamic response of the instrument and the sample, the following expressions are derived:

K − mω 2 =

Dω =

Fo cos φ ho

(1)

Fo sin φ ho

(2)

where m is the mass of the indenter column, ω the angular excitation frequency, Fo is the excitation amplitude force, ho is the displacement response and φ is the phase angle between Fo and ho. D accounts for the damping in both indenter head, Di, and sample, Ds (D = Di + Ds). An equivalent stiffness K is related to the stiffness of the contact, S, the load-frame stiffness, Kf, and the stiffness of the support springs, Ks, through the following expression: −1

1 1  K = + + Ks S K  f  

(3)

Combining equations 1 and 3 yields: 7



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    1 −1   S= −Kf  Fo  2 cos φ − ( − ω ) K m   s  ho 

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−1

(4)

In practice, m and Ks are determined from the measurement of the indenter hanging in free space and Kf is determined following the procedure described in reference 14. Assuming elastic-viscoelastic correspondence, modulus and hardness can be determined using:22,23

π 1 E = S 2 2 β Ac 1 −ν

H=

(5)

P Ac

(6)

Here, ν is Poisson’s ratio, taken to be 0.35 in all cases, β is a geometric factor (β = 1.034 for a Berkovich indenter) and Ac is the contact area at an applied load P. The area function describing Ac was calculated as a function of the contact penetration depth, hc, using a fused silica standard. On the other hand, hc was estimated using hc = h - 0.75 P/S where h is the total penetration depth. 14

3. RESULTS

8


6 5

E [GPa]

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


4 3 2 1 0 400

H [MPa]

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300 200 100 0 0

1000

2000

3000

4000

5000

Displacement, h [nm]

Figure 1. Elastic modulus E and hardness H variation as a function of displacement into the surface for iPP (), PPS (), iPP/IF-WS2 8% ( ) and PPS/IF-WS2 8% ( ).

3.1. Variation of the Mechanical Properties with Indentation Depth Figure 1 illustrates the modulus and hardness variation as a function of displacement into the surface for the neat iPP and PPS and the nanocomposites with the highest fullerene-like content (8 wt%). The H and E values plotted in this figure result from the average of at least 10 different DSI tests, the error bars representing the standard deviations. It is seen that at small penetration depths, H and E exhibit large deviations most probably due to surface roughness. These errors are substantially minimized in the case of the iPP nanocomposite with 8 wt% of filler where one can clearly distinguish quite constant average H and E values starting at a very small indentation depth (h ≈ 100 nm). In the rest of materials, the constancy of H and E with penetration depth is only distinct above h ≈ 1 µm where the error bars are significantly

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E [GPa]

5

4

3

2 400

H [MPa]

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300 200 180 160 140 0

1

2

3

4

5

6

7

8

% IF-WS2

Figure 2. E and H as function of filler content for the two types of nanocomposites ). Solid lines are eye-guides. investigated, iPP/IF-WS2 () and PPS/IF-WS2 (

shorter. Hence, for all the materials investigated, there is no indication of a change in mechanical properties as the indenter probe progresses across the thickness, from the surface towards the bulk, suggesting that the nanocomposites are homogeneous with respect to matrix morphology and filler dispersion. Figure 1 additionally shows that the average E and H values at deep penetrations significantly grow when inorganic fullerene particles are added to either PPS or iPP.

3.2. Mechanical Enhancement with Filler Content

Figure 2 illustrates the variation of

E and H as a function of filler content for the two series of composites based on iPP and PPS.

10


0.4

(E-Eneat)/Eneat

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0.2

0.0

-0.2 0.6

(H-Hneat)/Hneat

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0.4

0.2

0.0 0

1

2

3

4

5

6

7

8

% IF-WS2

Figure 3. Relative increase of E and H of each nanocomposite, (E-Eneat)/Eneat and (H-Hneat)/Hneat respectively, with respect to the value of the corresponding neat polymer (Eneat and Hneat). Symbols and lines as in Figure 2.

Values of E and H for each sample are calculated as the average within the interval of displacement into the surface: h = 3000 - 4500 nm (see Figure 1). For both nanocomposites, a clear E and H-rise is observed with increasing filler content up to ≈ 0.5 wt%, however, taking a closer look some differences can be observed. First of all, E and H of the PPS composites start to rise from the lowest loading, while in case of iPP there is first an initial drop. Most interesting is the observation that the reinforcement seems to be more efficient when added to the PPS matrix. Indeed, Figure 3 (and Table 1) collects the relative increase of E and H of each composite sample with respect to the corresponding initial material, defined as (E-

Eneat)/Eneat and (H-Hneat)/Hneat, where Eneat and Hneat are the modulus and hardness of the neat 11



0.7

ωc

0.6

0.5

0.4

D110 [nm]

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20

15

10 0

1

2

3

4

5

6

7

8

% IF-WS2

Figure 4. Degree of crystallinity, ωc, and average lateral crystal size in the (110) direction, D110, as a function of filler content for the iPP/IF-WS2 and PPS/IF-WS2 nanocomposites. Symbols and lines as in Figure 2.

polymer, respectively. It is noteworthy that at all filler contents, the mechanical properties of PPS are enhanced to a larger extent than the homologous ones of iPP. This difference is most conspicuous for H with respect to E and, in both cases, for high filler loadings: for instance, the addition of 8 wt% of IF-WS2 to iPP induces an enhancement for H and E of 18% each, whereas the same filler concentration in the PPS matrix yields H and E increases of 61% and 35% respectively. In order to further investigate the origin of the different reinforcement effect of the fullerene-like nanoparticles in the two matrices studied, next section includes results on the changes induced in the nanostructure of both polymers upon addition of the filler.

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3.3. Influence of the Filler on the Crystallinity and Average Crystal Size of the Matrix

Figure 4 illustrates the plots of the degree of crystallinity measured by X-ray

diffraction, ωc, and the average crystal size in the [110] direction (perpendicular to the (110) planes), D110, as a function of filler content for the iPP/IF-WS2 and the PPS/IF-WS2 nanocomposites. The degree of crystallinity shows a sudden drop in the presence of the filler, followed by a continuous increase that ends up with a plateau at a content of nanoparticles ≥ 1% wt. At first sight it seems to follow a similar trend as those of H and E. However, some differences that will be further analyzed in the following section can be distinguished between the E, H and the ωc behavior. The most significant one is that the initial drop in the degree of crystallinity found for the lowest filler content in the PPS/IF-WS2 nanocomposite does not mimic the almost constant values of E and H (see Figure 2 and Table 1). In addition, the D110 values of Figure 4 reveal that, in case of PPS, the increase or decrease of lateral crystal size is concurrent to the increase or decrease of crystallinity; both ωc and D110 show an initial drop for the lowest filler followed by a continuous increase with increasing filler content up to the leveling-off at filler contents of 1% wt. A parallel behavior of D110 and ωc is also observed for iPP at the lowest filler contents (0.05-0.1 wt%). However, for higher filler loadings, ωc increases while D110 decreases up to the final leveling-off. This finding can be explained as due to the strong nucleating effect of IF-WS2 in the iPP matrix that results in an enhanced crystallization rate, reduced size of the spherulites and improved levels of crystallinity.7

4. DISCUSSION 4.1. Mechanical Properties - Nanostructure Correlations

Figure 5 illustrates

the plot of the E and H values of all the nanocomposites (empty symbols) and their corresponding neat polymers (big solid symbols) as a function of the degree of crystallinity. 13



6

E [GPa]

5 4 3 2 1 0 500 D 11

0

=1

8n

m

400 H [MPa]

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= 14

300

D110

200

= D110

nm

23 n

m

100 0 0.0

0.2

0.4

ωc

0.6

0.8

1.0

Figure 5. E and H as a function of ωc, for all the nanocomposites investigated and their corresponding neat polymers: iPP (), iPP/IF-WS2 (), PPS () and PPS/IF-WS2 (). Data for an amorphous PPS film (ωc = 0) and an additional PPS sample with higher crystallinity (ωc = 0.584) are also included. The arrow points at the H value of the PPS/IF-WS2 0.5% sample. The lines follow equation 7 (see text).

For the sake of the discussion, the data for two additional PPS films, one amorphous and the other with a higher degree of crystallinity (achieved using a slower cooling rate from the melt) are also included in Figure 5. The figure shows that for each nanocomposite series, E and H increase with the degree of crystallinity. Previous studies have provided vast experimental evidence of a direct proportionality between H and ωc through:31-33

H = Hc ωc + Ha (1−ωc)

(7) 14


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where Hc accounts for the hardness of the crystals and Ha for the contribution of the amorphous regions. With respect to E, a number of models have been developed to describe the dependence with ωc including the parallel model, the series one, or a combination of both.34,35 Such correlation has been little explored using modulus data extracted from depthsensing instrumentation and only some indications that E could follow a parallel model with

ωc, as H does, have been proposed. 35 Indeed, in the present investigation, the similarity of the E and H variation with ωc (see Figure 5) suggests that the behavior of both magnitudes could be discussed analogously. However, to avoid unnecessary repetitions, what follows will be just focused on H.

4.2.

Reinforcement

in

the

Rubbery

Amorphous

Phase:

iPP/IF-WS2

Nanocomposites The H linear dependence with crystallinity (equation 7) has been represented in Figure 5 for the neat iPP sample (solid line). Here it has been assumed Ha ≈ 0 because the Tg of this material is below the temperature of measurement. The intersection of the extrapolated line with the right-hand y-axis represents the Hc value for this particular neat iPP (≈ 250 MPa). It is now widely accepted that Hc depends on the crystal thickness.31 Increasing crystal thickness values would yield straight lines with increasing slope (and hence, higher right-hand y-axis intercepts, Hc). The interesting observation now is that all the iPP/IF-WS2 nanocomposites exhibit smaller lateral crystal sizes than the reference iPP material (see Figure 4), and hence, presumably, smaller crystal thicknesses. Hence, one would expect all the composites H data to lie beneath the straight solid line represented in Figure 5. However, this is not the case and the H data of the iPP/IF composites either lie in the straight line describing the hardness of the matrix or even slightly above. This result indicates that the hardness variation associated to the incorporation of the filler is not only due to modifications 15



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in the polymer matrix nanostructure but also to the reinforcement of the filler itself. We have attempted to separately evaluate these two contributions by adopting the following procedure. Firstly, the maximum hardness of iPP for each nanocomposite material, HiPP, should be defined as the value predicted by the straight line of Figure 5 (HiPP = Hc ωc). One can easily calculate these values using Hc = 250 MPa and the respective ωc values shown in Figure 4 (see also Table 1). Then it follows that the minimum hardness enhancement due to the reinforcement of the filler, ∆HIF, can be estimated from ∆HIF = H- HiPP. The HiPP and ∆HIF values calculated for each nanocomposite sample are included in Table 1. The figures in parenthesis represent the fraction of maximum increase that can be attributed to changes in nanostructure, (HiPP - Hneat)/Hneat and to the minimum variation due to the filler reinforcement, ∆HIF/ Hneat. Note that (HiPP - Hneat)/Hneat + ∆HIF/ Hneat = (H-Hneat)/Hneat. The HiPP and ∆HIF

data of Table 1 address important observations for the iPP

nanocomposites: i) addition of an IF nanofiller with intrinsic higher hardness than the matrix yields a H enhancement itself (∆HIF) that seems to contribute more significantly than that attributed to changes induced in the iPP morphology (HiPP); ii) low filler loadings (≤ 0.1%) do not seem to influence the mechanical properties except via the small changes generated in the matrix morphology; iii) for higher filler loadings (> 0.1%), the hardness enhancement associated to the distribution of ‘hard’ particles in a ‘softer’ matrix (∆HIF) seems to be independent of the filler content approaching values of ≈ 10% increase with respect to the neat polymer matrix.

4.3.

Reinforcement

in

the

Glassy

Amorphous

Regions:

PPS/IF-WS2

Nanocomposites Following the same procedure as above, a straight dashed line passing through the hardness value of the amorphous PPS film and that of the crystallized PPS sample describes the variation of the hardness of PPS with the degree of crystallinity for crystal sizes with lateral dimensions around D110 ≈ 14 nm. It is clearly seen that all the values for the

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PPS/IF-WS2 nanocomposites deviate from this dashed line. The question now is whether this deviation can be attributed to changes in the crystal size or is it due to the reinforcement effect of the filler. To answer this question, a new PPS sample with higher degree of crystallinity and larger crystal size was prepared (this can be easily done as explained in reference 9). The

H and E values of this second PPS sample with ωc = 0.584 and D110 = 18 nm are included in Figure 5. As expected, the data of this sample lie far beyond the dashed line describing the hardness of PPS with D110 ≈ 14 nm. Accordingly, a new dotted line describing the hardness of PPS samples with D110 ≈ 18 nm can be represented. The influence of the size of the crystals on H of PPS can now be clearly discerned. Let us now select one of the nanocomposite materials with D110 ≈ 14 nm. This is the case for 0.5 wt% IF-WS2 (marked with an arrow in Figure 5, see the D110 data in Figure 4). The hardness value associated to the PPS matrix within this nanocomposite should lie in the dashed line of Figure 5 with D110 ≈ 14 nm. Deviations from this line should be attributed to the reinforcement of the filler itself, ∆HIF, and not to changes in the polymer nanostructure. One can now easily estimate the hardness of PPS for the PPS/IF-WS2 (0.5% wt%) nanocomposite, HPPS, using equation 7 with Hc = 342 MPa, Ha = 134 MPa and ωc = 0.596. The HPPS and ∆HIF (∆HIF= H- HPPS) values are included in Table 1.

It is noteworthy that the remarkable H-increase (48%) detected with the

incorporation of only 0.5% of filler can be mostly attributed to the nanoparticles reinforcement (42%) and only a small 6% of this increase is a consequence of changes in the polymer nanostructure. For the rest of polymer nanocomposites, one can estimate HPPS and

∆HIF in the same manner as that suggested for PPS/IF-WS2 (0.5% wt%). However, the following remark must be considered: for samples with D110 smaller than that of the neat PPS material (IF = 0.05, 0.1 %), only the upper and lower limits to the HPPS and ∆HIF values respectively can be calculated; in the case that D110 is higher than that of pure PPS (IF ≥ 1%),

17



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Page 18 of 30

the procedure can only give an estimate of the minimum value of HPPS and the maximum one of ∆HIF. Table 1 collects the values of HPPS and ∆HIF for all the nanocomposites investigated. Comparison of the HPPS data of Table 1 with those of HiPP shows that the reinforcement of PPS attributed to changes in the polymer nanostructure is very similar to that found for iPP (typically less than 10% increase with respect to H of the neat polymer). However, fundamental differences are found between their ∆HIF contributions. First, in the case of the PPS nanocomposites, ∆HIF/Hneat is significantly larger than (HPPS - Hneat)/Hneat (see Figures in parenthesis, Table 1) indicating that the reinforcement effect of the harder IF-WS2 particles in the PPS matrix is the major contribution to the hardness enhancement. Notable values of ∆HIF are detected in case of PPS/IF-WS2 composites, even at very low filler loadings (= 18% for a filler content of only 0.1%), and approaching very significant values at the highest IF contents (≈ 40-50% for filler loadings in the range 0.5 – 8%). This behaviour is in contrast to that found for the iPP nanocomposites where the hardness enhancement attributed to ∆HIF is very limited. The result is possibly related to the fact that the amorphous phase within the iPP matrix exhibits a viscous character at the temperature of measurement, a few tens of degrees above Tg (≈ 0 ºC).7 Under these conditions, one could envisage a mechanism of deformation where the spherical nanoparticles dispersed mainly in the amorphous phase would ‘slide’ easily across the viscous amorphous phase under the influence of the stress field of the indenter. In contrast, for the PPS/IF-WS2 nanocomposites, the glassy nature of the amorphous regions (Tg ≈ 90 ºC)9 provides the necessary conditions for an adequate transference of the stress field underneath the indenter throughout the sample. The filler particles are in this case embedded in a solid phase and would experience a deformation process induced by the local stress field. In preceding sections, the larger difference found between the H enhancements of both matrices upon filler loading with respect to those found for E was highlighted (see Figure 3,

18


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6 5 EINDENT [GPa]

Page 19 of 30

4 3 2 1 0 0

1

2

3 4 E´DMA [GPa]

5

6

Figure 6. E values, obtained from depth-sensing indentation, versus storage modulus values E´, determined by means of DMA at two different frequencies (1 Hz and 10 Hz), for all the nanocomposites investigated: iPP at 1 Hz () and 10 Hz (); iPP/IF-WS2 at 1 Hz () and 10 Hz (); PPS at 1 Hz () and 10 Hz () and PPS/IF-WS2 at 1 Hz () and 10 Hz (). DMA data at 1 Hz are taken from references 7 and 9. In case of neat PPS, only the semicrystalline sample with the lowest degree of crystallinity is included. The error for neat iPP lies within the symbol size.

Table 1). This result can now be explained in view of the distinctive mechanisms of matrixparticle load transfer proposed above for iPP and PPS that should express their largest differences in the regime of plasticity, where irreversible deformations take place, and to a lesser extent in the elastic regime, where strains are recovered upon load release. Finally, it is worth pointing out that the H levelling-off detected at filler loadings of approximately 0.5 – 1 wt% (see Figure 2), for both PPS and iPP, seems to be associated to a parallel levelling-off of the ∆HIF values (see Table 1). Increasing the filler content above 1 wt% does not significantly enhance the properties. This result is possibly associated to an increase in the particle-particle interaction at high filler loadings that decreases the matrixparticle load transfer. Indeed, scanning electron microscopy studies on iPP/IF-WS2 19



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Page 20 of 30

nanocomposites suggest the appearance of few-particle clusters for filler concentrations above 2 wt.%.7

4.4. Correlation with Macroscopic DMA Measurements Figure 6 illustrates the plot of the elastic modulus values measured by means of DSI, E (values taken from Table 1), versus the storage modulus values determined by means of DMA, E´, at two different frequencies (1 and 10 Hz) for all the nanocomposites investigated. Data at 1 Hz were taken from references 7, 9. The dotted straight line of Figure 6 represents the identity function. It is clearly seen that the indentation data are significantly higher (by an approximate factor of ≈ 1.5) than the values obtained by means of DMA. Increasing one order of magnitude the frequency of measurement used in DMA analysis only slightly increases the E´ modulus values without a significant consequence in the E/E´ ratio. Hence, it seems that the different frequencies employed in the mechanical tests (45 Hz for CSM indentation; 1 and 10 Hz for DMA) does not seem to explain the distinct E and E´ values encountered. On the other hand, it has been suggested that indentation and DMA data are at variance because the former technique yields information at a local scale while the latter applies to macroscopic properties. However, in the present work, the CSM technique has been used to follow the evolution of E (and H) with penetration depth. The results of Figure 1 demonstrated that E (and H) is constant for h ≥ 1 µm and this E levelling-off value at deep penetrations have been employed in our following discussions. It is noteworthy that the field of elastic deformation extends beyond the plastic-elastic boundary, the latter usually assumed to initiate at ≈ 10 times the indentation depth. Taking into account that indentations of 5 µm have been produced in our polymer samples, then the field of elastic deformation should start at penetration depths of ≈ 50 µm and one must conclude that our indentation data are indeed representative of the mechanical bulk properties of the polymer samples tested.

20


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There is a fundamental difference between the indentation and the DMA techniques that is most likely to be the origin of the divergence of results; this is the direction of application of the stress. In a tensile DMA test, the applied force is unidirectional while in the indentation test the stress field is believed to evolve radially from the onset of indentation, i.e., in a triaxial manner. In the former technique, mainly tensile stresses are applied while in the latter a combination of compressive and shear forces are exerted on the material. It is noteworthy that E values derived from conventional compressive testing are usually found to exhibit significantly higher values than those determined by means of tensile testing.35 The fact that E values from DSI data are higher than those determined from DMA is in agreement with these findings. In summary, although a number of papers report successful comparisons between modulus values determined by means of DMA and indentation, to our opinion, there is no sound physical basis for such correlation.

5. CONCLUSIONS Depth-sensing instrumentation has been used to evaluate the effect of inorganic fullerene-like nanoparticles in two different polymer matrices, iPP and PPS. In both cases, a mechanical reinforcement effect has been detected: both, hardness and modulus, initially rise with increasing filler content and finally level-off at filler loadings of ≈ 0.5-1%. An effort has been done to independently evaluate the effect of two main contributions to the mechanical improvement: i) the changes in the polymer matrix nanostructure induced by addition of the filler and ii) the reinforcement due to the intrinsic properties of the IF-WS2. Analysis of the indentation hardness results suggests that the changes induced in the matrix upon addition of the filler represent no more than 10% of the hardness increase. It is found that the mechanical behavior observed as a function of filler loading is mainly 21



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Page 22 of 30

attributed to the intrinsic reinforcement effect of the nanoparticles. In this respect, important differences are found between the iPP and the PPS nanocomposites. In the case of iPP, the filler reinforcement is limited to values around 10% or below, while for the PPS series, remarkable hardness enhancements up to 50% are found. It is suggested that the nature of the amorphous phase plays an important role on the overall resistance of the polymer nanocomposite to plastic deformation. A different mechanism of deformation is suggested for the iPP and the PPS nanocomposites. In the former case, the effective reinforcement of the filler is thought to be attenuated by the surrounding viscous amorphous phase. Finally, comparison of the modulus values obtained by means of instrumented indentation with those obtained by means of DMA reveal that the former technique yields significantly higher values than the latter. The result is mainly attributed to the different directionality of the applied forces.

ACKNOWLEDGEMENTS The authors wish to thank the MICINN (Ministerio de Ciencia e Innovación), Spain, for funding the research reported under the grant FIS2010-18069 and MAT2010-21070C02-01. MN would like to acknowledge the MINECO for a ‘Ramón y Cajal’ Senior Research Fellowship and AD wishes to acknowledge the CSIC for a JAE Postdoctoral Fellowship.

REFERENCES (1) Tenne, R.; Margulis, L.; Genut, M.; Hodes, G. Polyhedral and Cylindrical Structures of Tungsten Disulfide. Nature 1992, 360, 444-446.

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1, 103-111. (4) Tenne, R.; Redlich, M. Recent Progress in the Research of Inorganic Fullerene-like Nanoparticles and Inorganic Nanotubes. Chem. Soc. Rev. 2010, 39, 1423-1434. (5) Naffakh, M.; Díez-Pascual, A. M.; Marco, C.; Ellis, G. J.; Gómez-Fatou, M. A. Opportunities and Challenges in the Use of Inorganic Fullerene-like Nanoparticles to Produce Advanced Polymer Nanocomposites. Prog. Polym. Sci. 2013, 38, 1163-1231. (6) Tevet, O.; Goldbart, O.; Cohen, S. R.; Rosentsveig, R.; Popovitz-Biro, R.; Wagner, H. D.; Tenne, R. Nanocompression of Individual Multilayered Polyhedral Nanoparticles.

Nanotechnology 2010; 21, 365705-365710. (7) Naffakh, M.; Martín, Z.; Fanegas, N.; Marco, C.; Gómez, M. A.; Jiménez, I. Influence of Inorganic Fullerene-like WS2 Nanoparticles on the Thermal Behavior of Isotactic Polypropylene. J. Polym. Sci. Part B Polym. Phys. 2007, 45, 2309–2321. (8) Naffakh, M.; Marco, C.; Gómez, M. A.; Jiménez, I. Unique Isothermal Crystallization Behavior of Novel Polyphenylene Sulfide/Inorganic Fullerene-like WS2 Nanocomposites. J.

Phys. Chem. B. 2008, 112, 14819-14828. (9) Naffakh, M.; Marco, C.; Gómez, M. A.; Gómez-Herrero, J; Jiménez, I. Use of Inorganic Fullerene-like

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to

Produce

New

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Polyphenylene

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Nanocomposites: Role of the Nanoparticle Concentration. J. Phys. Chem. B. 2009, 113, 10104-10111.

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(10) Naffakh, M.; Marco, C.; Gómez, M. A.; Jiménez, I. Unique Nucleation Activity of Inorganic Fullerene-like WS2 Nanoparticles in Polyphenylene Sulfide Nanocomposites: Isokinetic and Isoconversional Study of Dynamic Crystallization Kinetics J. Phys.

Chem. B. 2009, 113, 7107-7115. (11) Naffakh, M.; Díez-Pascual, A. M.; Marco, C.; Gómez, M. A.; Jiménez, I. Novel MeltProcessable Poly(ether ether ketone)(PEEK)/Inorganic Fullerene-like WS2 Nanoparticles for Critical Applications. J. Phys. Chem. B 2010, 114, 11444-11453. (12) Naffakh, M.; Marco, C.; Gómez, M. A.; Jiménez, I. Novel Melt-processable Nylon6/Inorganic Fullerene-like WS2 Nanocomposites for Critical Applications. Mater. Chem.

Phys. 2011, 129, 641-648. (13) Doerner, M. F.; Nix, W. D. A Method for Interpreting the Data from Depth-sensing Indentation Instruments. J. Mater. Res. 1986, 1(4), 601-609. (14) Oliver, W. C.; Pharr, G. M. An Improved Technique for Determining Hardness and Elastic Modulus Using Load and Displacement Sensing Indentation Experiments. J.

Mater. Res. 1992, 7(6), 1564-1583. (15) Briscoe, B. J.; Fiori, L.; Pelillo, E. Nano-indentation of Polymeric Surfaces. J.

Phys. D: Appl. Phys. 1998, 31, 2395–405. (16) Flores, A.; Baltá-Calleja F. J. Mechanical Properties of Poly(ethylene terephthalate) at the Near Surface from Depth-sensing Experiments. Phil. Mag. A 1998,

78(6), 1283-1297. (17) Hochstetter, G.; Jimenez, A.; Loubet, J. L. Strain-rate Effects on Hardness of Glassy Polymers in the Nanoscale Range. Comparison between Quasi-static and Continuous Stiffness Measurements. J. Macromol. Sci. Phys. 1999, 38(5&6), 681-692. 24


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(18) VanLandingham, M. R.; Villarubia, J. S.; Guthrie, W. F.; Meyers, G. F. Nanoindentation of Polymers: An Overview. Macromol. Symp. 2001, 167, 15-43. (19) Herbert, E. G.; Oliver, W. C.; Lumsdaine, A.; Pharr G. M. Measuring the Constitutive Behavior of Viscoelastic Solids in the Time and Frequency Domain using Flat Punch Nanoindentation J. Mater. Res. 2009, 24(3), 626-637. (20) Hayes, S. A.; Goruppa, A. A.; Jones, F. R. Dynamic Nanoindentation as a Tool for the Examination of Polymeric Materials. J. Mater. Res. 2004, 19(11), 3298-3306. (21) White, C. C.; VanLandingham, M. R.; Drzal, P. L.; Chang, N-K.; Chang, S-H. Viscoelastic Characterization of Polymers Using Instrumented Indentation. II. Dynamic Testing. J. Polym. Sci.: Part B: Polym. Phys. 2005, 43, 1812–1824. (22) Herbert, E. G.; Oliver, W. C.; Pharr, G. M. Nanoindentation and the Dynamic Characterization of Viscoelastic Solids. J. Phys. D: Appl. Phys. 2008, 41, 074021 (9pp). (23) Liu, T.; Phang, I. Y.; Shen, L.; Chow, S. Y.; Zhang, W. D. Morphology and Mechanical Properties of Multiwalled Carbon Nanotubes Reinforced Nylon-6 Composites. Macromolecules 2004, 37, 7214-7222. (24) Shen, L.; Phang, I. Y.; Chen, L. Liu, T.; Zeng, K. Nanoindentation and Morphological Studies on Nylon 66 Nanocomposites. I. Effect of Clay Loading.

Polymer 2004, 45, 3341–3349. (25) Dutta, A. K.; Penumadu, D.; Files, B. Nanoindentation Testing for Evaluating Modulus and Hardness of Single-walled Carbon Nanotube-reinforced Epoxy Composites. J. Mater. Res. 2004, 19(1), 158-164.

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Table 1. Modulus E, Hardness H, Relative Increases with Respect to the Neat Polymer (E-Eneat)/Eneat and (H-Hneat)/Hneat, Degree of or PPS ωc, Matrix Hardness HiPP and Hardness Increase due to Filler Reinforcement ∆HIF Matrix wt% IF-WS2 E [GPa] H [MPa] (E-Eneat)/Eneat a (H-Hneat)/Hneat b HiPP or PPS [MPa] d ∆HIF [MPa] ωc c

Crystallinity

iPP

PPS

a

e

0

2.7 ± 0.1

153 ± 12

0

0

0.604

0.05

2.41 ± 0.09

147 ± 7

-0.11

-0.04

0.582

< 147 (-0.04)

>0

(0.0)

0.1

2.7 ± 0.1

152 ± 9

0

-0.01

0.598

< 151 (-0.01)

>1

(0.0)

0.25

3.00 ± 0.08

175 ± 6

0.11

0.14

0.626

< 158 (0.03)

> 17 (0.11)

0.5

2.95 ± 0.08

172 ±7

0.09

0.12

0.619

< 157 (0.02)

> 15 (0.10)

1

2.93 ± 0.08

165 ± 9

0.09

0.08

0.634

< 160 (0.05)

> 5 (0.03)

2

3.1 ± 0.1

182 ±6

0.15

0.19

0.649

< 164 (0.07)

> 18 (0.12)

4

3.13 ± 0.05

175 ±7

0.16

0.14

0.641

< 162 (0.06)

> 13 (0.08)

8

3.19 ± 0.05

180 ± 4

0.18

0.18

0.657

< 166 (0.09)

> 14 (0.09)

0

3.7 ± 0.3

242± 31

0

0

0.519

0.05

3.7 ± 0.2

251 ± 15

0

0.04

0.473

< 232 (-0.04)

> 19 (0.08)

0.1

4.0 ± 0.1

285 ± 20

0.07

0.18

0.521

< 242 (0.0)

> 43 (0.18)

0.5

4.7 ± 0.2

359 ±26

0.26

0.48

0.596

258

(0.06)

101 (0.42)

1

4.5 ± 0.2

336 ± 37

0.20

0.39

0.582

> 255 (0.05)

< 81 (0.34)

2

4.95 ± 0.08

382 ± 15

0.34

0.58

0.607

> 260 (0.08)

< 122 (0.50)

4

4.8 ± 0.2

350 ± 30

0.29

0.45

0.603

> 259 (0.07)

< 91 (0.38)

8

5.0 ± 0.1

390 ± 20

0.35

0.61

0.619

> 263 (0.08)

< 127 (0.53)

Eneat is the modulus of the neat polymer.

28


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b

Hneat is the hardness of the neat polymer. The degree of crystallinity ωc is measured by X-ray diffraction. d HiPP or PPS is the hardness value predicted by equation 7 associated to the iPP or the PPS matrix. The figures in parenthesis represent the fraction of increase that can be attributed to changes in nanostructure, i.e., (HiPP or PPS – Hneat)/Hneat. e ∆HIF is the hardness enhancement due to the reinforcement of the filler. The figures in parenthesis represent ∆HIF/ Hm. Note that for each material the addition of the figures in parenthesis in the last two columns yields (H-Hneat)/Hneat. c

29



Evaluating the reinforcement of inorganic fullerene-like nanoparticles in thermoplastic matrices by depth-sensing indentation A. Flores, M. Naffakh, A. M. Díez-Pascual, F. Ania and M. A. Gómez-Fatou

Reinforcement due to intrinsic properties of filler

∆H IF/Hneat

Overall Hardness enhancement

PPS

0.6

0.6

PPS

0.3

iPP

0.0

0

2

4

6

8

% IF-WS2 iPP

0.0 0

2

4

% IF-WS2

6

8

matrix

-Hneat)/Hneat

0.3

(H

(H-Hneat)/Hneat

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48

Page 30 of 30

0.6

Matrix = PPS, iPP

0.3 0.0

0

2

4

6

8

% IF-WS2 Enhancement due to changes in the matrix nanostructure

for Table of Contents use only

30


Evaluating the Reinforcement of Inorganic Fullerene-like

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