Bhattarai et al. [55] investigated CS–PCL polyblend nanofibres in comparison to PCL products. They reported an enhancement in mechanical properties and cellular behaviour without the creation of chemical crosslinking between the CS and the PCL. Copper et al. [15] also achieved better mechanical properties and cellular compatibility from CS-PCL fibres compared to the PCL products alone.

Furthermore, the different profiles of Figures 1 and 2 reveal that the fibre strength increased, but the IFSS decreased with increasing CS concentration. Therefore, an optimal solution can be expected between PCP-3R and -3 to balance the mechanical and interfacial properties during the application of the CS coupling agent.

#### *4.3. SEM Analysis*

As shown in Figure 3, the surface of the uncoated fibre (PCP-0) was smooth, while excess attachment was observed on the surface of all the coated fibres (PCP-3R, -3, -6, and -9), which evidenced the presence of a CS coating on the fibres. The difference between Figure 3b,c indicated removal of excess coating by post-cleaning. In addition, no visible cracks or pits were observed on the surface of all the fibres after dipping or post-cleaning treatments.

#### *Fibers* **2018**, *6*, 97

The blistering and skinning indicated by white arrows in Figure 3f were observed at the fibre surface after a prolonged exposure to the electron beam. This phenomenon was suggested to be associated with the heating on the fibre surface caused by electronic scanning, indicating a heat-sensitive coating layer on the fibre. This was interpreted as further evidence of a polymeric CS layer on the fibres. Furthermore, the connection between two filaments presented in in Figure 3g showed the potential of CS to act as a binder or film former as well as a coupling agent.

#### *4.4. TGA*

Zawadzki et al. [27] conducted TGA of CS from room temperature to 600 ◦C under a nitrogen atmosphere, and suggested that the TGA behaviour of solid CS can be divided into three stages: (1) the removal of water below 100 ◦C, in which the water was physically adsorbed and/or weakly hydrogen-bonded to CS molecules; (2) the release of strongly hydrogen-bonded water starting above 100 ◦C and reaching the maximum rate at ~168 ◦C; and (3) the predominant weight loss by 43% at 230–400 ◦C caused by the depolymerisation of CS chains, the decomposition of pyranose rings via dehydration and deamination, and finally, a ring-opening reaction. The quantities of weight loss provided in Table 3 indicated the presence of a coating layer on the fibres. The thermal behaviours traced in Figure 4 showed a good agreement with reference [27], further proving the coating layer to be CS.

An increase in weight loss of CS was also shown in Table 3, indicating a thicker CS coating layer with higher CS concentration, which was suggested to be due to the aggregation behaviour of CS polymer chains. The rapid release of alkali (Na+) in the coating solution introduced negative charge at the fibre surface, which attracted NH3 + groups of CS, further leading to an uneven charge distribution of the CS chains caused by asymmetric electron dispersion. Then, the CS coating on the fibre surface was increased by charge attraction and stabilised via intermolecular hydrogen-bonding [27]. Several articles [10,56] have reported that CS can form an entangled network in an acidic aqueous medium by crosslinking with itself and that the crosslinking involves two structural units that may be from the same or from different CS polymer chains.

In addition, the post-cleaning process may have disrupted the intermolecular hydrogen bonds and consequently removed any excess CS, leaving only the firmly bonded CS on the fibre surface [25].

#### *4.5. FTIR and Raman Analyses*

As shown in Figure 5, the absorption bands at 3291 and 3358 cm–1 were assigned to the stretching vibrations of O–H and N–H, respectively. The intensity of these two bands on CS–S was observed to reduce as compared to CS–P, which was suggested to be due to the formation of intermolecular hydrogen bonds in the excess CS coating [11]. The band of NH2 at 1589 cm–1 was replaced by a new band of NH3 + at 1560 cm–1 on CS–S, suggesting the formation of hydrogen bonds to balance the change of charges in the excess CS [27]. The bands at 561 and 1063 cm–1 could be identified as crystallisation-sensitive bands of CS, and the changes of these bands in wavenumber and intensity that occurred on CS–S implied that an intermolecular interaction between different polymer chains disturbed the crystallisation of CS [34]. The absence of a characteristic band at ~1730 cm−1 on CS–S implied that there was no residual COOH group on the fibres after coating application [57].

The band at 495 cm−1 shown in Figure 6a was assigned to the asymmetric bending vibration of O–P–O in Q1 units, and a slight shift to a lower wavenumber occurred for the coated fibres, which might be attributed to the interaction between the pyrophosphate units and the coating [28]. The bands at 754 and 894 cm−1 were assigned to the symmetric and asymmetric stretching of P–O–P bridging bonds in Q2 units, respectively [30]. The latter band was observed to decrease in intensity and shift to a lower wavenumber for the coated fibres, which was suggested to be due to the reduction in Q2 units, indicating a depolymerisation of metaphosphate units in acid solution. The Raman spectra in Figure 7a also showed a reduction in intensity of the band at 1097 cm−1 (the asymmetric stretching vibration of (P2O7) 4− in Q1 units [38]) and the shoulder at ~1250 cm−1 (the asymmetric stretching

vibration of (PO3) − in Q2 units [39]) for the coated fibres, indicating a decrease in concentration of pyro- and metaphosphate units.

As shown in Figure 6b, the replacement of the band of NH2 at 1589 cm–1 by the new bands appearing at 1547, 1560, and 1552 cm–1 on PCP-3, -6, and -9, respectively, was attributed to the protonation of the amine groups in CS [33]. Lawrie et al. [58] reported that partial protonation of NH2 groups in CS could cause the appearance of a new band at ~1530 cm−1 due to one of the NH3 + vibrational modes. The difference of the band locations between the literature and this study might be due to the interaction between the charged amine groups and the phosphate units. The Raman spectra in Figure 7b,c show that the bands of the bending (1596 cm−1) and stretching (3305 cm−1) vibrations of NH2 disappeared for the coated fibres, and this was attributed to the protonation of amine groups, which correlated well with the IR results [41].

Bhumkar et al. [59] investigated the crosslinking of CS with sodium tripolyphosphate in acidic aqueous solution via FTIR. They reported that in the IR spectra of crosslinked CS, the band at 1655 cm–1 disappeared and two new bands appeared at 1554 and 1645 cm–1, due to the linkage between the phosphoric and ammonium ions. Therefore, the appearance of the bands at ~1547 and ~1635 cm−1 and the intensity reduction of the band at 1648 cm−1 for the coated fibres suggested interaction of the NH3 + groups with the fibre surface (see Figure 6b). Moreover, it was also suggested that the formation of hydrogen bonding contributed to the interaction between the charged CS and PGFs. Sayyar et al. [60] characterised graphene-filled CS composites using FTIR. They found that the bands corresponding to the N–H bending of the amino group (1573 cm−1), C=O stretching of the amide group (1658 cm−1), and N–H stretching of the amino group (3464 cm−1) in CS shifted to a lower wavenumber in composite films, indicating likely hydrogen-bonding interactions between CS and lactic acid and reduced graphene oxide.

#### *4.6. XPS Analysis*

The C1s peaks of CS shown in Figure 8a have a good agreement with the literature [61]. The peak at 284.5 eV was assigned to C–C or C–NH2; 286.0 eV was assigned to C–O, C–OH, or C–NHC=O; and 287.6 eV was assigned to O–C–O or C=O–NHR chemical bonds. The detection of these signals for the coated fibres indicated the presence of CS on the fibre surface. The minor difference between the spectra of the CS powder and the coated fibres was suggested to be due to contamination carbon on the fibres.

The N1s peak of CS could be resolved into two peaks at 399.1 and 400.7 eV, assigned to the nitrogen of the amino and amide groups, respectively [62]. The atom % of the amino peak was 95%, agreeing well with the degree of deacetylation (DD > 80%) quoted by the supplier. As shown in Figure 8b, the peak of NH2 was also found for the coated fibres, whilst a new peak of NH3 + appeared at 401.3 ± 0.1 eV binding energy, indicating the partial protonation of the amino groups [62]. Lawrie et al. [58] performed XPS on CS powders from Sigma with a reported 85% degree of deacetylation. They identified the N1s peak as one main signal at 399.4 eV with another weaker signal at 400.5 eV, corresponding to the amino and amide, respectively. They subsequently characterised the acid-treated CS film and detected one more signal at 401.4 eV, attributing to the protonated amine.

As shown in Table 5, the atomic percentage of the NH3 + in PCP-3 increased from 40% to 65% after the post-cleaning, supporting the assertion that the post-cleaning process removed excess unreacted CS adsorbed to fibre surface by weak hydrogen bonds. In contrast, the atom % decreased to 17% in PCP-6 with increasing CS concentration, indicating an increasing layer of excess CS. The atomic % was observed to remain constant in PCP-9 as compared to PCP-6, indicating that there was a threshold of concentration of CS interacting with the PGF surface. The changes in the atomic % of NH3 + were also observed to be consistent with the IFSS profile shown in Figure 2, which revealed the efficiency of the CS coating.

The O1s spectrum of CS was resolved into one main peak at 532.3 eV and two weak peaks at 530.9 and 533.7 eV. The peak at 532.3 eV was assigned to C–O–C, O–H, or bound water and the two weak peaks were assigned to C=O in the amide group and O–C–O in the pyranose ring, respectively [61]. The O1s spectrum of PGF was deconvoluted into two peaks at 530.6 and 531.8 eV binding energy, corresponding to the nonbridging and bridging phosphate oxygens, respectively [25]. Although it was difficult to differentiate the subpeaks from the overlapped binding energy range, the signal of CS was seen to dominate the overlapping peak gradually from PCP-3R to PCP-9, which indicated the increase in CS on the fibres, fitting well with the TGA results.

The P2p peak of PGF was composed of the P2p3/2/P2p1/2 doublet, and was fitted with an energy difference of 1.1 eV and an approximate atomic ratio of 2:1. The peak of P2p3/2 at 133.3 eV was attributed to the pentavalent tetracoordinated phosphorus (pyrophosphate and orthophosphate) surrounded by a different chemical environment (phosphate-like structure), and the peak of P2p1/2 at 134.4 eV was attributed to the metaphosphate [63]. The peaks at similar binding energy were also observed for the coated fibres.

In summary: (1) The presence of CS coating on the fibre surface was confirmed via TGA, FTIR, Raman, and XPS analyses. (2) SEM and TGA indicated the formation of a thicker CS coating layer when using the solution of higher CS concentration. (3) The results of tensile tests indicated that the CS coating protected the fibres from losing strength in the acid solution. (4) The results from single-fibre fragmentation tests indicated that CS coupling agents improved the IFSS of the composites, due to the interaction between the protonated amine groups of CS and the hydrogen groups in the fibre surface, as correlated via the analyses of FTIR, Raman, and XPS.

#### **5. Conclusions**

CS showed its potential to improve the interface of PGF/PCL composites, in which the protonation of amine groups plays a central role in the interaction between CS and PGF. The nitrogen spectrum of XPS revealed a threshold of efficient CS coating for interfacial improvement. The post-cleaning process could remove the excess CS on the fibre surface and consequently increase the efficiency of the coupling agent.

The coated CS layer showed a significant effect on the fibre surface protection, resulting in a maintenance of mechanical strength after dipping in the acetic acid solution. In addition, TGA and SEM indicated an increase in coating thickness with higher CS concentration in the coating solution. However, a thicker CS layer could also have led to poorer load transmission between the fibre and matrix. Therefore, a balance of the mechanical and interfacial properties would be considered during the application of CS as a coupling agent. This balance is suggested to be achieved by adjusting the CS concentration and the post-cleaning duration.

**Author Contributions:** C.T. and I.A. conceived the original idea. C.T. carried out the experiment. J.Z. fabricated the PCL film and single fibre composite samples. W.C. supported on the XPS experiment and discussion. C.T. wrote the manuscript with support from I.A., A.P. and N.S. A.P. helped supervise the project. C.R. and I.A. supervised the project.

**Funding:** The authors kindly acknowledge the financial support from the Ningbo Education Bureau, Ningbo Science and Technology Bureau, China's MOST, and The University of Nottingham. Additionally, the authors also would like to acknowledge financial support from the Ningbo S&T bureau Ningbo International Collaboration Project (project code 2017D10012) and the Ningbo 3315 Innovation team Scheme "Composites Development and Manufacturing for Sustainable Environment".

**Acknowledgments:** The authors would like to acknowledge all the technical staff of the ITRC and Wolfson Building in UK campus and the SEB in Ningbo campus of Nottingham University, for their guidance and technical assistance.

**Conflicts of Interest:** The authors declare no conflicts of interest.

#### **References**

1. Ahmed, I.; Jones, I.A.; Parsons, A.J.; Bernard, J.; Farmer, J.; Scotchford, C.A.; Walker, G.S.; Rudd, C.D. Composites for bone repair: Phosphate glass fibre reinforced PLA with varying fibre architecture. *J. Mater. Sci. Mater. Med.* **2011**, *22*, 1825–1834. [CrossRef]

© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Further Progress in Functional Interlayers with Controlled Mechanical Properties Designed for Glass Fiber/Polyester Composites**

#### **Antonin Knob 1, Jaroslav Lukes 2, Lawrence Thadeus Drzal 3 and Vladimir Cech 4,\***

Received: 19 July 2018; Accepted: 14 August 2018; Published: 16 August 2018

**Abstract:** Compatible interlayers must be coated on reinforcing fibers to ensure effective stress transfer from the polymer matrix to the fiber in high-performance polymer composites. The mechanical properties of the interlayer, and its interfacial adhesion on both interfaces with the fiber and polymer matrix are among the key parameters that control the performance of polymer composite through the interphase region. Plasma-synthesized interlayers, in the form of variable materials from polymer-like to glass-like films with a Young's modulus of 10–52 GPa, were deposited on unsized glass fibers used as reinforcements in glass fiber/polyester composites. Modulus Mapping (dynamic nanoindentation testing) was successfully used to examine the mechanical properties across the interphase region on cross-sections of the model composite in order to distinguish the fiber, the interlayer, and the modified and bulk polymer matrix. The interfacial shear strength for plasma-coated fibers in glass fiber/polyester composites, determined from the microindentation test, was up to 36% higher than those of commercially sized fibers. The effects of fiber pretreatment, single and double interlayers, and post-treatment of the interlayer on interfacial shear strength were also discussed. Functional interlayers with high shear yield strength and controlled physicochemical properties are promising for high-performance polymer composites with a controlled interphase.

**Keywords:** glass fiber; polymer-matrix composites; interface/interphase; nanoindentation; mechanical properties; microindentation test

#### **1. Introduction**

Fiber-reinforced polymer composites are significant materials for transportation, building/construction, electrical/electronic applications, and consumer goods mainly because of their low density and high specific strength. The performance of the polymer composite is controlled by the properties of the fiber reinforcement, the polymer matrix, and the interphase [1], which is the region of material between the fiber and the polymer matrix responsible for the effective stress transfer from the matrix to the fiber. The composite interphase comprises a functional interlayer, and eventually, a modified matrix region. The functional interlayer is a material in the form of thin film that is deposited on the fiber surface to improve compatibility and create a strong but tenacious link between both components. The modified matrix is the part of the polymer matrix affected by the presence of the coated fiber. It is the interphase that is very important for the resultant performance of

the composite material. Even reinforcing fibers with high modulus or high strength and polymers with high thermal or chemical resistance result in inefficient composites in the case of a poor interphase. The interphase is built in polymer composites using functional interlayers through fiber sizing, coating, or grafting in a wet chemical process or through chemical vapor deposition [2].

Investigating interphase properties is very important for understanding its functionality and the possibility of its optimization for a specific composite system, i.e., the fiber and the polymer matrix. The reinforcement and the polymer matrix are mostly very different materials in terms of their mechanical, electrical, and chemical properties. The interphase can, therefore, be identified between the two composite phases based on differences in material structure, Young's modulus, hardness, coefficient of friction, dielectric permittivity, elemental composition, or chemical structure. The cross-section of the polymer composite was investigated using microscopic infrared spectroscopy [3]; however, the interphase region is tiny since the interphase thickness is generally less than several microns, and nanoscale imaging techniques are necessary to ensure sufficient lateral resolution. Transmission electron microscopy (TEM) has sufficient lateral resolution to distinguish the interphase region based on differences in the atomic mass of the elements, and even the element composition profile across the composite phases can be determined using electron energy loss spectroscopy (EELS) [4,5]. The results show some limitations depending on the technique used to fabricate a sample for TEM imaging, such as focus ion beam (FIB), ion beam etching, and ultramicrotomy [6].

Another principal technique suitable for interphase characterization is atomic force microscopy (AFM), which is used in different modes. Standard AFM tips with a radius of less than 10 nm provide sufficient lateral resolution. Electrostatic force microscopy (EFM) is one of the modes used to measure electrostatic interactions between the AFM tip and the composite phases made from dielectric materials. The interaction is sensitive to charge distribution or local differences in dielectric permittivity of composite phases [7]. During scanning in phase-imaging AFM mode, the oscillating tip is influenced by contact with composite components of different viscoelastic properties, resulting in a phase shift of the tip against the applied oscillating force. The interphase area can be identified due to changes in phase distribution over the sample surface [8,9]. Similarly, lateral force microscopy using AFM contact mode can be used to characterize the distribution of lateral forces on the sample surface to distinguish the interphase area due to a different coefficient of friction [9]. In this mode, the constant force-feedback on the sample adjusts the height of the AFM cantilever to compensate for the topographic features of the surface, and the torsion of the cantilever bearing the tip characterizes lateral forces as a result of surface friction. In the intermodulation AFM, the cantilever is simultaneously excited by two frequencies near the resonance, and the cantilever response at a new frequency contains information on the surface mechanical properties of the composite sample [10]. Atomic force acoustic microscopy (AFAM) is a combination of ultrasonics with AFM, where the composite sample is excited with ultrasonic frequency by means of a piezoelectric transducer that causes vibrations of the sample surface, and the vibration behavior of the AFM cantilever in contact with the surface of the sample is sensitive to its local elastic properties [9,11]. All of the aforementioned AFM modes have some limitations in order to sufficiently differentiate the interphase region.

The difference between the mechanical properties of composite constituents is sufficient to potentially distinguish the interphase region. Nanoindentation techniques [12,13] are, therefore, ideal for characterizing the distribution of mechanical properties across the fiber/matrix interface in terms of Young's modulus and hardness. Quantitative nanomechanical measurements are possible, starting with forces ≤1 μN, but the lateral resolution is not as high as in AFM measurements due to a higher radius of commercially available diamond indenters with a radius of 50 nm (Berkovich) or 40 nm (cube corner). However, nanoindentation together with nanoscratch measurements is influenced by fiber constraint [8], because the indenter induces stress near the fiber surface, and the critical distance is approximately twenty times the indentation depth [14,15]. Dynamic nanoindentation may be operated as a Modulus Mapping mode that appears to be more suitable for observing the cross-section of the

composite sample using detailed maps of the complex modulus, the storage, and the loss modulus, just a few nanometers below the sample surface, thereby avoiding the fiber constraint [9,16]. This technique is carried out by superimposing a small sinusoidal force on a greater but constant force during the indentation test.

Recently, force mapping spectroscopy (nanomechanical mapping) [17–21] was applied in order to characterize the interphase region in polymer composites and nanocomposites. In this AFM mode, a force–distance curve is recorded in each pixel of the map, and a linear portion of this curve corresponding to an elastic regime is used to determine the Young's modulus using the Derjaguin–Muller–Toporov (DMT) model. However, the determined values of Young's modulus are not realistic for the range of moduli from the polymer matrix to the reinforcement, and they are strongly affected by the surface topography of the composite sample. It was found that surface formations (grains), with a radius similar or larger than that of the AFM tip or diamond indenter, have a critical influence on probe (AFM or nanoindentation) measurements used to characterize mechanical properties based on the geometry of contact between the tip and the sample [22].

Plasma-enhanced chemical vapor deposition (PECVD) is a technique for depositing thin films with controlled mechanical properties [23], which are intended to be used as functional interlayers for glass fiber/polyester composites in order to improve their interfacial shear strength (IFSS). This study is focused on examining the mechanical properties of such interlayers in model composites using Modulus Mapping based on our previous experience [9]. Selected interlayers with controlled mechanical and chemical properties were tested to determine how they affect the IFSS in glass fiber/polyester composites determined by the microindentation test [24]. Also, the effects of pretreatment, plasma coating (thin-film deposition), and post-treatment of glass fibers (GF) on determined IFSS were firstly discussed.

#### **2. Materials and Methods**

#### *2.1. Thin-Film Deposition*

A plasma reactor [25] equipped with a radio-frequency (RF; 13.56 MHz) capacitive coupling system using asymmetric plan-parallel plate electrodes was employed to deposit thin films on silicon wafers (polished on both sides (1 0 0) with an overlayer of native SiO2 (3 nm), infrared (IR), transparent, 0.8 × 10 × 10 mm3; ON Semiconductor, Roznov pod Radhostem, Czech Republic) and single GFs (unsized (bare) E-glass, 1200 tex, mean diameter 19 μm; Saint-Gobain Adfors, Litomysl Czech Republic). Fifteen GFs separated from the GF bundle were placed parallel to each other on silicon wafers and fixed with adhesive carbon tape suitable for vacuum applications. Hydrogenated amorphous carbon-silicon (a-CSi:H) films were deposited from pure tetravinylsilane (TVS; Si–(CH=CH2)4 purity 97%, Sigma Aldrich, Prague, Czech Republic) precursor, and hydrogenated amorphous carbon-silicon oxide (a-CSiO:H) films were deposited from TVS in a mixture with oxygen gas (TVS/O2 mixture) using PECVD. Firstly, the silicon wafer (alone or with single GFs) was loaded into the RF bottom electrode, and the plasma reactor was evacuated to a basic pressure of 1 × 10−5 Pa. Then, the silicon wafer was pretreated with argon (Ar) plasma (10.0 sccm, 5.7 Pa, continuous wave 5.0 W) for 10 min to remove adsorbed gases and to ensure the reproducible adhesion of deposited films. The a-CSi:H or a-CSiO:H films were deposited using pulsed plasma at a mass flow rate of 3.8 sccm, a pressure of 2.7 Pa, and an effective power of 2.0–10 W. The effective power, *W*eff, was set by changing the ratio of the time the plasma is switched on, *t*on, to the period, *T*, defined as *T* = *t*on + *t*off, and thus, *W*eff = *t*on/*T* × *W*total, where *t*off is the time the plasma is switched off and *W*total is the total RF power. An oxygen fraction of 0.92 was used for the TVS/O2 mixture as the mass flow rates for TVS and O2 gas were 0.30 sccm and 3.5 sccm, respectively. A movable shutter was used to deposit the films in steady-state plasma conditions monitored by mass spectrometry (Process Gas Analyser HPR-30, Hiden Analytical, Warrington, UK). When the deposition process was completed, the plasma reactor was flushed with Ar gas (10.0 sccm, 5.7 Pa) to remove residual gases for 60 min, and the reactor was

then evacuated to a basic pressure of 1 × 10−5 Pa. The deposited sample was moved to the load-lock chamber after 12 h, and the chamber was then flooded with air to atmospheric pressure [25].

Another plasma reactor [26], an RF helical coupling system operated in a pulsed regime, was used to deposit (PECVD) thin films on special glass slides without flaws (1.0 × 26 × 76 mm3; Knittel Glaser, Braunschweig, Germany), silicon wafers, and GF bundles specified above. This cylindrical discharge is axially symmetrical, and therefore, more suitable for film deposition on the GF bundle mounted in the discharge axis. This plasma reactor was first evacuated to a basic pressure of 5 × 10−4 Pa. The planar substrate or GF bundle was pretreated with O2 plasma (5.0 sccm, 4.0 Pa, 25 W) for 60 min to desorb gases and remove hydrocarbons from the substrate surface (planar or fibrous). The oxygen fraction X=O2/(TVS + O2) in the gas mixture was varied from 0–0.46 (0, 0.10, 0.21, 0.33, and 0.46) at a total flow rate (TVS + O2) of 0.80 sccm, and the corresponding pressure was 1.4 Pa. To deposit all films, an effective power of 5.0 W (time-on pulse, period, and total power were 1 ms, 10 ms, and 50 W, respectively) was used. Finally, after film deposition, the apparatus was flushed with Ar gas (10.0 sccm, 10 Pa). After 60 min, the reactor was evacuated to a basic pressure, and, after 12 h, was flooded with air at atmospheric pressure.

#### *2.2. Thin-Film Analysis*

The thickness of the films deposited on the silicon wafers was determined using a phase-modulated spectroscopic ellipsometer, UVISEL (HORIBA Scientific, Longjumeau, France), at a wavelength of 250–830 nm in a 2-nm step; the angle of incidence was 70◦ and the spot size was 100 × 300 μm2. The average deposition rate was given by the ratio of film thickness to deposition time. The thickness of the film deposited on planar or fibrous substrate was controlled by the deposition time at a constant deposition rate.

Infrared measurements in the 500–4000 cm−1 wavenumber range were made using a VERTEX 80 vacuum Fourier-transform infrared (FTIR) spectrometer (Bruker Optics, Billerica, MA, USA). Transmission spectra were obtained on films deposited on infrared-transparent silicon wafers. To remove the spectral properties of the silicon wafer, an absorption-subtraction technique was used, and background correction was carried out before each measurement. The spectral resolution was 4 cm−1, and 256 scans were recorded in order to achieve a reasonable signal-to-noise ratio. The bulk chemical composition of the thin films was investigated using conventional and resonant Rutherford backscattering spectrometry (RBS) and elastic recoil detection analysis (ERDA) methods on a Tandetron 4130MC accelerator (High Voltage Engineering Europa B.V., Amersfoort, The Netherlands (HVEE)) using 2.73-MeV alpha particles.

The Young's modulus of the films deposited on the silicon wafers was determined with the Oliver–Pharr method [27] from nanoindentation measurements using a two-dimensional (2D) TriboScope (Hysitron, Eden Prairie, MN, USA) attached to an NTegra Prima Scanning Probe Microscope (NT-MDT), Zelenograd, Russia. A Berkovich tip with a radius of curvature of 50 nm was used. The experiment was performed in cyclic mode to obtain the depth profile of Young's modulus to 15–20% of the film thickness, applying 19 loading and partial unloading cycles in a single indentation, while the unloading fraction was kept constant at 0.8, and the loading was increased using a displacement exponent of 2 (exponential function) [28]. The experimental depth profile was extrapolated to the zero-contact depth (film surface) to estimate the correct Young's modulus of the film not influenced by the substrate properties [29].

#### *2.3. Sample Preparation and Modulus Mapping*

After thin-film deposition, 15 GFs fixed with adhesive carbon tape were transferred from the silicon wafer into a silicon rubber mold, which was filled with unsaturated polyester resin (isophthalic) Viapal HP 349 F (Sirca S. p. A., Sandono di Massanzago, Italy), and cured at 140 ◦C to form a polymer block, 5 × 13 × 25 mm3 with unidirectional fibers. This block was inserted into the metallographic

sample holder with GFs placed normally to the sample surface, and the surface was polished using conventional metallographic techniques.

The mechanical properties of the interphase region around a single GF in the polyester matrix were characterized using a Hysitron TI 950 TriboIndenter nanomechanical test instrument (Hysitron, Eden Prairie, MN, USA) with a nanoDMA III (Dynamic Mechanical Analysis) package. The Modulus Mapping mode combines the in situ scanning probe microscopy (SPM) imaging capability of Hysitron's nanomechanical testing instruments with the ability to perform nanoDMA III tests. During the test, the indenter is scanned over the surface to image the topography of the sample. During scanning, the probe is oscillated with the specified frequency and amplitude of the load. The resulting signal from the transducer is then analyzed to determine the displacement amplitude and phase delay of the oscillation at each pixel in the image [9]. The software plots these measurements in separate image files, and after scanning is complete, the images can be analyzed to determine the storage, *E*', and loss, *E*", modulus at each point using the following equations:

$$E' = \frac{k\_s \sqrt{\pi}}{2\sqrt{A}}, \ E'' = \frac{\omega \mathbb{C}\_s \sqrt{\pi}}{2\sqrt{A}} \tag{1}$$

where *k*s and *C*s are the stiffness and damping of the specimen determined from the displacement amplitude and phase difference between the force and the displacement, *A* is the projected contact area, and ω is the frequency used [27,30]. A cube diamond indenter (NorthStar, Hysitron, Eden Prairie, MN, USA) with a radius of curvature of 40 nm was used for Modulus Mapping. The amplitude of a dynamic force of 5 μN at an oscillation frequency of 220 Hz was superimposed to a normal contact force of 15 μN in order to maintain the dynamic displacement amplitude at 4 nm when scanning the sample surface in an in situ SPM mode. The radius of the tip was carefully calibrated using Modulus Mapping in standard silica at the above conditions, and with the same displacement as in the test specimen. The map resolution in the 3 × 3 μm2 scan area was 256 × 256 pixels.

#### *2.4. Composite Fabrication and Microindentation Test*

After the GF bundle was surface-modified in an RF helical coupling system, the bundle was impregnated with unsaturated polyester resin and then axially placed in a silicon rubber mold that was filled with resin, before being cured at 140 ◦C to form a polymer disc of 14 mm in diameter and 5 mm in height [31]. Similar to the resin block with unidirectional fibers, the disc was embedded in a metallographic specimen and polished using conventional metallographic techniques.

The polished cross-section of the composite, with the fibers normal to the surface of the sample 3 mm high, was used to characterize the interfacial adhesion of GF in the polyester matrix using the microindentation test [24]. Using an Interfacial Testing System (ITS) (Dow Chemical Company, Midland, MI, USA) [32], a diamond tip with a diameter of 12 μm was applied to push the single fiber from its surrounding polyester resin. The tested fiber was indented at increased normal loading until failure of adhesion. The de-bonding of the fiber is visible as a dark shadow around the fiber. When the shadow appears around the fiber at an angle of between 90◦ and 120◦, the fiber is considered to be de-bonded [32]. The IFSS was determined from the de-bond load, *P*, using a generalized empirical equation,

$$IFSS = A \frac{P}{D^2} \left[ B \left( \frac{G\_m}{E\_f} \right)^{\frac{1}{2}} - C \log \left( \frac{d}{D} \right) - E \right] \tag{2}$$

where *D* is the fiber diameter, *Gm* and *Ef* are the matrix shear modulus and fiber axial modulus, respectively, *d* is the matrix thickness between the tested fiber and its nearest neighbor, and *A* = 1.249 × 104, *B* = 0.8757, *C* = 0.01863, and *E* = 0.02,650. A matrix shear modulus of 1.5 GPa and a fiber axial modulus of 73 GPa correspond to the polyester resin used and the GF, respectively [32].

#### **3. Results**

#### *3.1. Mapping of Storage Modulus across the Interphase Region*

Two polymer blocks, 5 × 13 × 25 mm3, each with 15 unidirectional fibers (mean diameter 19 μm) embedded in polyester resin, were fabricated using unsized and commercially sized GFs to examine the interphase region with Modulus Mapping. The commercial sizing based on silane coupling agents was designated by the GF manufacturer for GF/polyester composites. After polishing, the metallographic specimen with GFs placed normally to the surface was scanned in a 3 × 3 μm2 area using a cube corner indenter with a radius of 40 nm. The storage modulus across the fiber/polymer interface for the unsized GF is shown in Figure 1a. A three-dimensional (3D) image with a storage modulus in a color scale, which is used for better orientation, shows the black area corresponding to the polymer matrix and a part of the fiber with a sharply raised modulus at the fiber/polymer interface. The 2D projection of the storage modulus is visible at the top of Figure 1a. A similar image of the storage modulus was obtained for the commercially sized GF (Figure 1b). The detailed profile of the storage modulus close to the fiber/polymer interface (Figure 1c) enabled the finding of characteristic differences. The red line for the unsized GF rises in the immediate vicinity of the fiber surface, and the 0.07-μm-thick area on which the modulus was changed corresponds to the fiber constraint area rather than the interphase region. However, the blue line corresponding to the commercially sized GF indicates the wider area of the modulus change that can be assigned to the interphase region with a thickness of 0.5 μm. The interphase region around the commercially sized GF was non-regular, and the thickness varied from 0.1 (GF without sizing) to 0.5 μm. The storage modulus for the polyester resin outside the interphase region was approximately 6 GPa, as shown in Figure 1c, which appears to be a little overvalued because a Young's modulus of 4 GPa was evaluated from the quasi-static tensile test of the polyester resin. An increased storage modulus was found in the subsurface layer up to a depth of 40 nm of metallographic specimen, probably due to resin polishing.

**Figure 1.** Storage modulus mapping of the interphase region in model glass fiber (GF)/polyester composites for (**a**) unsized GF, (**b**) sized GF, and (**c**) a detailed profile close to the fiber/polymer interface.

The polymer-like a-CSi:H film was deposited (68.7 nm/min) on single GFs and silicon wafers at an effective power of 2.0 W (time-on pulse, period, and total power were 1 ms, 5 ms, and 10 W, respectively) from a pure TVS precursor using the plasma reactor with plan-parallel electrodes. The film thickness was 895 nm, which was determined from the ellipsometric spectrum of the film on the silicon wafer, and we expect that approximately the same thickness of the film was deposited on the GFs. A Young's modulus of 10.0 ± 0.5 GPa was evaluated for the film on the silicon wafer using cyclic nanoindentation measurements. RBS/ERDA spectra of a-CSi:H film allowed determining the chemical composition and calculating the element ratios to the silicon concentration, resulting in a stoichiometric formula a-C5.6Si:H5.0. The high concentration of hydrogen (43 atom%) that is bonded to silicon and carbon atoms is responsible for the relatively low cross-linking of the mainly carbon network in the a-CSi:H film, resulting in the low Young's modulus that corresponds to a polymer-like material. GFs coated with this interlayer (a-C5.6Si:H5.0 film) were inserted in polyester resin to fabricate a model composite, whose cross-section was analyzed with Modulus Mapping to distinguish the interlayer between the fiber and the polymer matrix, and to characterize its mechanical properties. The storage modulus in a color scale across the interphase region in a model GF/polyester composite is shown in Figure 2a. The lateral resolution of this map is 11.7 nm. The scan area of 3 × 3 μm2 shows the white area with a modulus above 62 GPa corresponding to GF surrounded by a ring with a modulus similar to the polymer matrix (black area on the left) with a blue contour. Profiles of the storage modulus (red line) and surface topography (blue line), corresponding to the GF with the polymer-like interlayer, are plotted in Figure 2b. Interpretation of mechanical properties is not easy due to the surface topography of the composite sample marked with a blue dashed line. Generally, the indentation measurement analysis is based on the assumption that the process is initiated by penetrating the indenter into a flat surface. The influence of surface roughness on nanoindentation measurements was interpreted using model simulations [22]. The polymer composite is a hybrid material, and its polished surface exhibits roughness in the interphase region due to the different mechanical properties of the composite phases. Based on the thickness of the interlayer, its position is expected at the GF surface, indicated by the yellow area in Figure 2b. At the interface between the interlayer and the polymer matrix, there is an elevated modulus, but it is an artefact due to the side contact of the indenter with the rough surface of the composite sample. In this case, the measured contact stiffness is higher than that corresponding to the flat surface, and the projected contact area is underestimated, resulting in an overestimated storage modulus. The slope of the surface across the interlayer provides a slightly underestimated modulus, corresponding to model simulations [22], but part of the interlayer close to the interlayer/fiber interface corresponds to the correct value of the storage modulus of 10 GPa. The surface roughness of the interphase region complicates the differentiation between the polymer-like interlayer and the polymer matrix with similar mechanical properties.

**Figure 2.** Modulus Mapping (scan area 3 × 3 μm2) across the interphase region with a polymer-like (hydrogenated amorphous carbon-silicon (a-C5.6Si:H5.0)) interlayer in a model GF/polyester composite: (**a**) map of the storage modulus, (**b**) storage modulus and topography profiles.

An interlayer with a higher Young's modulus of 16.3 ± 0.8 GPa was deposited (9.7 nm/min) on single GFs at an effective power of 2.0 W (1 ms, 5 ms, and 10 W) from TVS in a mixture with oxygen gas with a 0.92 oxygen fraction. Oxygen atoms (34 atom%) were incorporated into the thin film, but carbon (29 atom%) and hydrogen (24 atom%) concentrations were reduced in the a-C2.2SiO2.6:H1.8 film with a thickness of 966 nm, unlike the polymer-like a-C5.6Si:H5.0 film. Cross-linking of the carbon-silicon oxide network is still relatively low due to Si–H, Si–OH, C–H, C–OH, and C=O bonding species. The results of Modulus Mapping across the interphase region in a model GF/polyester composite with a-C2.2SiO2.6:H1.8 interlayer are plotted in Figure 3. The map of the storage modulus in Figure 3a provides easy orientation in specific composite phases; the red/white area corresponds to a fiber, surrounded by a blue belt, indicating an interlayer with a higher modulus than the polymer matrix that is marked with a black area. The surface topography across the interlayer, marked by the yellow area, was flat, and was, therefore, favorable for determining the correct value of the storage modulus (Figure 3b). The storage modulus of the interlayer slightly increased from the interface between the interlayer and the polymer matrix to the interlayer/fiber interface, but the mean value was 16 GPa. A small portion of the interphase region with a thickness of 0.1–0.2 μm between the interlayer and the polymer matrix had a slightly elevated modulus, and could, therefore, correspond to the modified matrix.

**Figure 3.** Modulus Mapping (scan area 3 × 3 μm2) across the interphase region with the a-C2.2SiO2.6:H1.8 interlayer in a model GF/polyester composite: (**a**) map of the storage modulus, (**b**) storage modulus and topography profiles.

Another interlayer in the form of a silicon oxide film with a carbon concentration of 16 atom% was deposited (7.5 nm/min) at 4.5 W (1 ms, 11 ms, and 50 W) from TVS in a mixture with oxygen gas with a 0.92 oxygen fraction. The a-C0.7SiO2.1:H0.8 interlayer with a thickness of 1109 nm was more cross-linked, with a Young's modulus of 29.6 ± 1.5 GPa. Modulus Mapping in Figure 4a distinguished the fiber (white area) from the interlayer (purple area) and the polymer matrix (black area). A part of the blue belt with a thickness of 0.3 to 0.4 μm between the interlayer and the polymer matrix may represent a modified matrix region characterized by a gradually increasing modulus. The beginning of the step in the surface topography between the modified and the bulk matrix was responsible for the elevated modulus, which is an artefact caused by a specific surface topography, and can be seen as a purple contour in the blue belt (Figure 4b). In this case, the mechanical properties of the interlayer were sufficiently high to avoid any surface roughness between the fiber and the interlayer. This means that the surface topography was flat for the interlayer, delineated with the yellow area, and the storage modulus of the interlayer increased from 21 to 31 GPa at the interlayer/fiber interface at a mean value of 26 GPa.

**Figure 4.** Modulus Mapping (scan area 3 × 3 μm2) across the interphase region with a hydrogenated amorphous carbon-silicon oxide (a-C0.7SiO2.1:H0.8) interlayer in a model GF/polyester composite: (**a**) map of storage modulus, (**b**) storage modulus and topography profiles.

An SiO2-like interlayer with a small concentration of carbon (3 atom%) and hydrogen (11 atom%) was deposited (5.5 nm/min) at 10 W (1 ms, 5 ms, and 50 W) from TVS in a mixture with oxygen gas with a 0.92 oxygen fraction. The a-C0.1SiO1.9:H0.4 interlayer with a thickness of 950 nm was highly cross-linked resulting in a Young's modulus of 51.7 ± 3.3 GPa. Figure 5a shows a map of the storage modulus corresponding to the fiber (white area), interlayer (red area), modified matrix (blue area), and the bulk matrix (black area). The same deposition time resulted in consistent film thickness for films deposited on planar and fibrous substrates. The thickness of the modified matrix was increased for the SiO2-like interlayer, and varied from 0.5 to 0.7 μm. The mean value of the storage modulus for the flat surface of the interlayer was 52 GPa (Figure 5b), as expected from cyclic nanoindentation. The storage modulus across the interlayer varied like across the GF surface. This means that the interlayer material was as brittle as the glass, and the composite surface was specifically modified (slightly scratched) with 50-nm alumina nanoparticles, which were used to polish composite samples. The root-mean-square (RMS) roughness of the composite surface was 30 nm for a scan area of 100 × 100 μm2 measured by AFM. The varying modulus was caused by surface topography. Nanoindentation measurements of the material with a higher modulus are more sensitive to surface roughness [22]. The wide region of the modified matrix with the elevated modulus between the interlayer and the bulk matrix is evident for the SiO2-like interlayer in Figure 5b. We can assume that the modified matrix region was created as a result of a chemical bonding at the interface between the interlayer and the polymer matrix, unlike the unsized GF (Figure 1c). The polymer network was firmly anchored onto the surface of the interlayer, which increased the rigidity of the polymer network. The interlayer/polymer matrix interface was sharp, and no interpenetrating network between the interlayer and the polymer matrix can be expected, unlike commercially sized GFs that are surface-modified by silane coupling agents [33]. Analysis of the depth profile for the interlayers and GFs revealed that the subsurface layer, up to a depth of 5–10 nm, showed an increase in the storage modulus due to the mechanical modification of the material during the polishing procedure. Therefore, the storage modulus in the maps may be overestimated by up to 10% for the SiO2-like interlayer and GFs, as the average displacement of the indenter was approximately 5 nm, but it cannot affect the polymer-like interlayer with an average displacement of 15 nm. The differences in storage modulus among the composite phases were significantly higher compared to the differences in loss modulus, and were, therefore, more appropriate for the distinction of the specific phase.

**Figure 5.** Modulus Mapping (scan area 3 × 3 μm2) across the interphase region with the SiO2-like (a-C0.1SiO1.9:H0.4) interlayer in a model GF/polyester composite: (**a**) map of storage modulus, (**b**) storage modulus and topography profiles.

#### *3.2. Selected Interlayers for GF/Polyester Composite*

As shown in Section 3.1, a wide range of materials from polymer-like to SiO2-like films with a Young's modulus of 10 to 52 GPa can be used as interlayers for GF/polyester composites. The chemical and physical compatibility of the interlayer and the polymer matrix was increased for the polymer-like interlayers with the lower (≤16 GPa) Young's modulus as the thickness of the modified matrix region was decreased. Model simulations (finite element analysis) enabled calculating the interfacial stress fields in a GF/polyester composite during the microindentation test. The simulation showed that inserting a suitable interlayer between the fiber and the polymer matrix resulted in a significant reduction in shear stress across the interphase region when the polymer composite was under mechanical or thermal loading. If the indenter pushed the individual fiber into the polyester resin during the microindentation test, the interfacial shear stress at the interlayer/fiber interface increased with the raised Young's modulus of the interlayer [34]. The interfacial shear stress as a function of the indenter displacement for the interlayer Young's modulus in the range of 1 to 40 GPa is shown in Figure 6a. As can be seen from Figure 6a, the use of the interlayer with the higher modulus means that the interlayer/fiber interface must be stronger in order to avoid interphase shear failure. Also, the shear yield strength of the interlayer must be higher than the applied shear stress. Thermosetting or thermoplastic coatings are poor in shear [35], and are, therefore, unsuitable for use as functional interlayers in high-performance fiber-reinforced polymer composites. However, the plasma coatings deposited (PECVD) from organosilicon precursors have a sufficiently high shear yield strength at a relatively low Young's modulus [25], as opposed to thermosetting or thermoplastic coatings (see the schematic comparison in Figure 6b). Therefore, the plasma coatings may be more suitable as functional interlayers.

*Fibers* **2018**, *6*, 58

**Figure 6.** Model simulations: (**a**) shear stress at interlayer/fiber interface as a function of indenter displacement for interlayers with different Young's moduli during the microindentation test, (**b**) schematic correlation between shear yield strength and Young's modulus for plasma coatings [25] and thermosetting or thermoplastic coatings.

In our previous work, an oxygen-free TVS was successfully used for plasma coating of GFs in the RF helical coupling system to enhance the interfacial adhesion in GF/polyester composites [36]. In that case, the Young's modulus of the interlayer was increased from 9.4 to 23 GPa, reducing the vinyl concentration with an enhanced power of 0.1–5.0 W, but the adhesion of the interlayer to the GF did not significantly improve. In another study, therefore, a mixture of TVS precursor in gaseous oxygen was used to deposit plasma coatings onto GFs. Incorporating oxygen atoms into plasma coatings can increase interfacial adhesion, not only on the GF surface, but also at the interface with the polyester matrix due to better wettability and higher density of interfacial bonds. Indeed, the IFSS for oxidized plasma coatings was up to 21% higher than that for the non-oxidized interlayer, indicating the direct chemical effect of oxygen atoms on the interphase properties [34]. A nanoscratch test was used to measure the interlayer adhesion on a flat glass substrate, as well as on a single GF. For a given film, consistent interlayer adhesion was identified for both fibrous and planar glass substrates [37]. The work of adhesion was used as the right parameter for the characterization of the interlayer adhesion measured by a nanoscratch test. A strong correlation was found between the IFSS in GF/polyester composites and the interlayer adhesion for the same plasma coating [38]. Model and experimental data showed that the shear strength at the interlayer/fiber interface is the key factor affecting the mechanical response of the plasma-coated fibers in a GF/polyester composite at microindentation testing.

Further enhancement of the adhesion of oxidized plasma coatings deposited at higher power was our motivation to test an improvement of IFSS for GF/polyester composites. A set of interlayers was deposited onto silicon wafers and GF bundles at a constant effective power of 5.0 W, but the various oxygen fractions in the TVS/O2 mixture varied from 0 to 0.46 (0, 0.10, 0.21, 0.33, and 0.46) in the RF helical coupling system. Using the same power, the power density in the RF helical coupling system was significantly lower, due to a higher volume of plasma discharge, compared to the plasma reactor equipped with parallel electrodes. This means that plasma coatings synthetized in the RF helical coupling system contain a higher concentration of hydrogen atoms and are less cross-linked. The Young's modulus of plasma coatings deposited on silicon wafers was determined from nanoindentation measurements for various oxygen fractions in the TVS/O2 mixture (Figure 7a). The Young's modulus decreased from 22 to 18 GPa as a result of an increase in oxygen concentration from 0 to 11 atom% in a-CSiO:H films (Figure 7b). Chemical analysis (RBS/ERDA) showed that the C/Si and H/Si ratios in all films were approximately 7.1 and 10, respectively, but the O/Si ratio increased from 0 to 2.4 with the enhanced oxygen fraction (Figure 7b). FTIR spectra indicate that oxygen atoms were incorporated into hydroxyl, carbonyl, and Si–O–C/Si–O–Si species, and the intensity

and area of the absorption bands increased with enhanced oxygen fraction (Figure 7c). This trend is consistent with the previous analysis [39]. This means that the number of Si–C species in the a-CSi:H film deposited from pure TVS were reduced at the expense of the newly formed Si–O–C species in a-CSiO:H films synthetized from TVS in a mixture with oxygen gas. An increase in the number of Si–O–C/Si–O–Si species was important for increased adhesion at the interlayer/glass interface due to chemical reactions of plasma species with hydroxyl groups on the surface of the glass. By increasing the oxygen fraction, only a slight reduction of the vinyl groups in the film can be expected (Figure 7c). The vinyl groups on the interlayer surface are responsible for chemical bonding with the polyester resin during the curing process. The formation of hydroxyl and carbonyl groups eliminates one and two bonds, respectively, from each carbon atom that could potentially be used for cross-linking the network. Therefore, decreased cross-linking of the carbon-silicon oxide network with an increased oxygen fraction resulted in a film density reduction of 1.47 to 1.31 g/cm3 (Figure 7d), as determined from the RBS spectra, and also resulted in a reduced Young's modulus (Figure 7a).

**Figure 7.** Analysis of thin films deposited on silicon wafers at 5.0 W from the tetravinylsilane (TVS)/O2 mixture at different oxygen fractions 0–0.46: (**a**) Young's modulus determined from nanoindentation measurements, (**b**) oxygen concentration together with the O/Si ratio evaluated from Rutherford backscattering spectrometry (RBS)/elastic recoil detection analysis (ERDA) spectra, (**c**) infrared spectra for oxygen fractions 0–0.46, and (**d**) density of the film estimated from RBS spectra.

Interlayers with the properties shown in Figure 7 were also deposited on an unsized GF bundle of 1600 single filaments. The GF bundle was positioned along the axis of the axially symmetric plasma, resulting in the deposition of a uniform and homogeneous coating around the fibers [36]. TVS molecules were fragmented during the plasma process to form free radicals, and these highly reactive radicals recombined on the surface of the GFs to form an interlayer. Because of the low pressure (1.4 Pa) during the

plasma process, free radicals diffuse into the central part of the GF bundle and form an interlayer even on the surface of the central fibers. The thickness of the interlayer is controlled by deposition time at a known deposition rate, which is constant during deposition. However, the interlayer on the central fibers is thinner than that on the fibers at the edge of the bundle, because the deposition rate decreases radially into the GF bundle due to the shadow effect of the surrounding fibers. We found that the thickness of the film deposited on the silicon wafer under a single GF was reduced by a factor of 0.9, due to the shadow effect, after the deposition process with plan-parallel electrodes. Multiple shielding can be expressed in a geometric series as *tn* = *t*s *qn* − 1, where *t*s is the film thickness on the fiber at the edge of the bundle, *tn* is the film thickness on the *n*-th fiber in the direction toward the center of the bundle shielded by *n* − 1 fibers, and *q* is the shielding factor. For the GF bundle with a circular cross-section of 1600 single filaments, 22 is the order of the fiber in the center of bundle, and for *t*s = 1.0 μm, *q* = 0.9, the film thickness is *t*22 = 109 nm on the central GF. The thickness of the interlayer on the central GF is, therefore, ten times smaller than the thickness of the interlayer on the fiber at the edge of the bundle. The expected thickness of the interlayer varied between 1.0 and 0.1 μm across the GF bundle for all tested GFs. The plasma-coated GF bundle was embedded in the polyester resin, and was cured to form a unidirectional GF/polyester composite. The polished cross-section of the composite sample was subjected to microindentation testing for the IFSS evaluation. The mean value and its standard deviation were calculated from measurements of ten fibers selected from different parts of the bundle cross-section. No evident differences between IFSS values for central and edge fibers were observed, which corresponds to the previous results [34]. The dependence of IFSS on the oxygen fraction in the TVS/O2 mixture for plasma-coated GFs with an interlayer thickness of 1.0 μm (edge fibers) at an effective power of 5.0 W is given in Figure 8a. The IFSS (blue square) ranged from 130 to 139 MPa at a fiber volume fraction of approximately 0.53 in the GF/polyester composite. The maximum IFSS at a 0.33 oxygen fraction corresponds to the highest adhesion of the interlayer on glass substrate characterized by the work of adhesion (red circle) [38]. A 3.4-fold increase in the adhesion of the a-CSiO:H film (0.33 oxygen fraction), due to a raised concentration of Si–O–C/Si–O–Si bonding species at the interlayer/glass interface, compared to the a-CSi:H film (zero oxygen fraction) was promising for using the adhesive film as a functional interlayer in the GF/polyester composite. The IFSS trend corresponds to the adhesion trend (Figure 8a), but the statistical significance of the maximum IFSS is not entirely convincing. Although there are differences between the microindentation test and the nanoscratch test in the specimen loading and failure mechanics, and both the IFSS and the work of adhesion are calculated based on simplified assumptions, the interlayer adhesion on the glass surface was shown to control the shear strength in the GF/polyester composite [38].

**Figure 8.** Interfacial shear strength (IFSS) for plasma-coated glass GFs in the GF/polyester composite: (**a**) IFSS as a function of oxygen fraction in the TVS/O2 mixture used to deposit the interlayer at 5.0 W, (**b**) effect of GF pretreatment and post-treatment of plasma-coated GFs on IFSS for the same Interlayer A.

The effect of GF pretreatment and post-treatment of plasma-coated GFs on IFSS is demonstrated for the same Interlayer A (type A: 0.33 oxygen fraction, 5.0 W) in Figure 8b. The strong adhesion at the interlayer/fiber interface was primarily achieved by pretreatment of GFs with O2 plasma, which cleans and activates the fiber surface due to physical and chemical etching before the interlayer is deposited. The deposition of Interlayer A on GFs without pretreatment resulted in a significant reduction of IFSS to 83 MPa compared to 139 MPa for GFs pretreated with O2 plasma at 25 W for 60 min (Figure 8b). Ar gas (4.0 sccm) was mixed with O2 gas (0.5 sccm) to increase the degree of ionization during the pretreatment of GFs, but the IFSS of 139 GPa was the same as in the pretreatment with O2 plasma using the same Interlayer A. The IFSS value was reduced to 123 GPa when Ar/O2 plasma was applied for 45 min followed by O2 plasma for 15 min during the pretreatment procedure. This means that the cleaning and activation of the GF surface was not so effective, probably because of the interruption of the pretreatment process to change the gas mixture. This pretreatment method, however, was used for the other two samples, where, after the deposition of Interlayer A, the plasma-coated GFs were left in the reactor filled with TVS vapor for 10 s or 60 min without RF discharge. During this post-treatment procedure, the TVS molecules are not only physisorbed, but also chemisorbed on the chemically active surface of the interlayer that contains free radicals. An ultrathin overlayer, rich in vinyl groups, is formed on the interlayer surface during the application of TVS vapor. This layer may be important for improved adhesion at the polymer matrix/interlayer interface due to chemical bonding of vinyl groups with polyester resin during the curing process. Progress in the IFSS of 8 and 20 MPa (up to 143 MPa) was found for the interlayer, post-treated for 10 s and 60 min, respectively.

Interlayer B, deposited from pure TVS at an effective power of 0.1 W, contains no polar groups, but contains a higher concentration of vinyl groups compared to Interlayer A, as evident from the FTIR spectra (Figure 9a). Interlayer B could, therefore, be used to improve chemical bonding between the polymer matrix and the interlayer, similar to the ultrathin overlayer deposited from TVS vapor during the post-treatment procedure. In addition, Interlayer B has a lower Young's modulus of 9.4 GPa [36], which is beneficial for decreasing shear stress across the interphase region (Figure 6a). Interlayer B was combined with Interlayer A to form a bilayer that was tested to improve the IFSS in GF/polyester composites. Interlayer A, containing 6.1 atom% of oxygen, was deposited on the GF bundle followed by the oxygen-free Interlayer B in three thickness variations: 0.9 μm (Interlayer A) + 0.1 μm (Interlayer B), 0.1 μm (Interlayer A) + 0.9 μm (Interlayer B), and 0.05 μm (Interlayer A) + 0.05 μm (Interlayer B). The resulting IFSS values were compared with the single Interlayer A in Figure 9b. All the results in Figure 9b correspond to GFs pretreated with Ar/O2 plasma for 45 min, followed by O2 plasma for 15 min. The IFSS for bilayers was slightly higher by 1 to 4 MPa, which was statistically insignificant. The bilayer thickness was only 0.1 μm on edge fibers for the third bilayer, meaning that the bilayer thickness was only 0.01 μm on the central fibers. The final bilayer was prepared from the C and D interlayers by diluting the deposition plasma by adding 1.0 sccm of Ar gas to increase the ionization degree, and the total flow rate was, therefore, 1.8 sccm. Interlayer C was deposited from a TVS/O2/Ar mixture with a 0.15 oxygen fraction and a 0.55 argon fraction at an effective power of 5.0 W, but Interlayer D was synthesized from a TVS/Ar mixture with a 0.55 argon fraction at 0.1 W. The deposition rates for this bilayer were not known, but a significantly reduced IFSS value of 47 MPa indicates that the bilayer thickness was not sufficient for a uniform fiber coating, probably <0.1 μm (edge fibers). This means that part of the GF surface was not covered by the interlayer.

**Figure 9.** (**a**) Infrared spectra of Interlayers A and B, (**b**) effect of the bilayer on IFSS for the same pretreatment of GFs.

The final comparison of IFSS values for the GF/polyester composite in Figure 10 includes unsized, O2-plasma pretreated but uncoated, commercially sized, and optimized plasma-coated GFs. No chemical bonding at the polymer matrix/fiber interface for unsized GFs with an IFSS of 29 MPa is expected. O2-plasma pretreatment of GFs without deposition of an interlayer resulted in improved interfacial adhesion at the polymer matrix/fiber interface, but the IFSS at 49–50 MPa was limited by the shear failure of the matrix, as demonstrated by model simulations [34]. The IFSS of 105 MPa for commercially sized GFs was twice as high compared to GFs pretreated with O2 plasma. However, the optimized Interlayer A deposited on O2- or Ar/O2-plasma pretreated GFs resulted in an IFSS of 139 MPa. A further increase in IFSS to 143 MPa was achieved by the subsequent post-treatment of Interlayer A in TVS vapor for 60 min.

**Figure 10.** Interfacial shear strength (IFSS) for the GF/polyester composite reinforced with unsized, O2-plasma pretreated but uncoated, commercially sized, and optimized plasma-coated GFs.

Plasma coatings of tailored physicochemical properties have potential applications in high-performance polymer composites with a controlled interphase. Plasma nanotechnology is a promising technique for the construction of gradient interlayers needed for the new concept of composites without interfaces [40].

#### **4. Conclusions**

A wide range of polymer-to-glass materials was plasma synthesized (PECVD) to control their mechanical properties with Young's moduli from 10 to 52 GPa. The oxygen-free (a-CSi:H) and oxygen-bound (a-CSiO:H) films of 1.0-μm thickness were deposited from pure TVS and TVS mixed with oxygen gas on planar and fibrous substrates. Cyclic nanoindentation was used to determine the Young's modulus for films deposited on planar substrates. Single GFs coated by the interlayer were embedded into the polyester resin to form a model composite. The polished cross-section of the model composite was used to distinguish the composite phases, due to their different mechanical properties, with Modulus Mapping. The maps of storage modulus allowed defining the interlayer between the fiber surface and the polymer matrix. The results showed that the mechanical properties, along with the deposition rates, are consistent for films on planar and fibrous substrates. Modulus Mapping revealed that the thickness of the modified matrix increased up to 0.7 μm with the raised Young's modulus at 52 GPa for the SiO2-like interlayer. We assumed that the modified matrix region was created due to chemical bonding at the polymer matrix/interlayer interface, unlike the unsized fiber, where no interphase region was observed. Thus, a strong interfacial adhesion between the polymer matrix and stiffer material (interlayer) may result in the formation of a modified matrix region with gradually increased Young's modulus. The thickness of the modified matrix increased with the increasing difference between the interlayer and the polymer matrix moduli. Plasma organosilicon coatings with a relatively low Young's modulus (polymer-like material) have a shear yield strength that is significantly higher than that for thermosetting and thermoplastic coatings. This means that the plasma coatings with controlled interfacial adhesion and Young's modulus are promising as interlayers in polymer composites. Selected compatible interlayers with increased interlayer adhesion were deposited onto a 1600-filament GF bundle. The interlayer thickness on the fibers at the edge of the bundle was found to be reduced by one order for fibers in the center of the bundle as a result of the shadow effect. Optimized interlayer adhesion (0.33 oxygen fraction, 5.0 W) resulted in an IFSS of 139 MPa, supported by O2 and Ar/O2 pretreatment of GFs. The IFSS was reduced to 83 MPa if the pretreatment procedure was omitted. The GF pretreatment was, therefore, very important for cleaning and activating the surface of the fibers. Following the plasma coating, the optimized interlayer was then post-treated in TVS vapor to increase the concentration of vinyl groups on the interlayer surface. The post-treatment procedure increased the IFSS to 143 MPa, which is 3.9 times higher than unsized GFs, 1.9 times higher than O2-plasma pretreated GFs, and 36% higher than commercially sized GFs.

**Author Contributions:** A.K. deposited the interlayers, produced the composite samples, performed the microindentation tests, and analyzed the data; J.L. designed and performed the Modulus Mapping, and analyzed the data; L.T.D. provided the model simulations, as well as advice and discussions about the microindentation test and interface/interphase phenomena; V.C. provided the idea for this study, proposed the deposition conditions and experiments, analyzed and discussed the data, and wrote the paper.

**Funding:** This research was funded by the Grantová Agentura Cesk ˇ é Republiky, grant no. 16-09161S.

**Acknowledgments:** The authors would like to thank Petr Kucharcik for technical assistance, Erik Palesch for cyclic nanoindentation measurements, Vratislav Perina for RBS/ERDA analysis, and Miroslav Sirovy and Saint-Gobain Adfors CZ s.r.o. for providing glass fibers.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Investigation of Transcrystalline Interphases in Polypropylene/Glass Fiber Composites Using Micromechanical Tests**

#### **Hanna Brodowsky \* and Edith Mäder**

Leibniz-Institut für Polymerforschung Dresden, e.V. (IPF) Hohe Str. 6, D-01069 Dresden, Germany<EMAIL_ADDRESS>
**\*** Correspondence<EMAIL_ADDRESS>Tel.: +49-351-4658-641

Received: 20 December 2017; Accepted: 30 January 2018; Published: 12 March 2018

**Abstract:** In composites, a strong interphase between the components is essential for mechanical properties. By using a suitable sizing (i.e., surface modification) of the fiber, the interphase may be varied, e.g., by suppressing or promoting heterogeneous nucleation of a thermoplastic matrix. In the latter case, three-dimensional transcrystallized interphases with properties differing from those of the bulk matrix are formed. Polypropylene-glass fiber composites are prepared as single-fiber model composites with (a) sizings either inducing or suppressing a transcrystalline interphase, (b) different amounts of modifier maleic acid anhydride grafted polypropylene, and (c) different molecular weights of the matrix polymer. These are studied in quasi-static or cyclic load tests. Static tests permit insights in the interfacial characteristics such as critical interface energy release rate, adhesion strength and frictional stress. Cyclic tests on these model composites can be used to study the nature of dissipative processes and the damage behavior. Atomic Force Microscopy (AFM) investigations of the fiber fracture surfaces provide supplementary information. The transcrystalline layer can indeed improve the mechanical parameters (a 70–100% increase of strength and a 25 or 125% increase in toughness, depending on the molecular weight (MW) of the matrix polymer at low modifier concentration). However, the effect is partially neutralized by an opposing effect: high nucleation in the bulk in samples with commonly used concentrations of modifier.

**Keywords:** polypropylene; glass fiber; polymer-matrix composites; interface; mechanical behavior; transcrystallinity; micromechanical tests

#### **1. Introduction**

Semicrystalline thermoplastics crystallize in spherulites. In absence of external nuclei, the homogeneous crystallization can only occur at temperatures below the crystallization temperature Tcr. The nucleation rate is material-dependent, it generally increases with the temperature difference Tcr-T [1]. Once nucleated, the growth of spherulites in the bulk occurs in all directions until it reaches the neighboring spherulites, which impede further growth. On surfaces, the nucleation may occur by heterogeneous nucleation at the surface and can be significantly higher than in the bulk. In this case, the impedation by neighboring crystallites occurs almost immediately in the lateral directions, leaving only the outward direction for crystal growth. Thus, a so-called transcrystalline (TC) layer at the surface is formed. Transcrystallization is a nucleation controlled crystallization process which can occur in a semicrystalline polymer in contact with a second material [2,3]. The growth mechanism itself, as well as the growth rate, is identical in the bulk and within the TC layer [4–7]. The thickness of the TC layer is therefore determined by the different rates of the two nucleation processes in the bulk and at the surface, in relation to the growth rate [5,8,9]. These three quantities depend on a number

of thermodynamic and physicochemical conditions, such as the surface free energy, the nucleation density and the sample temperature.

In the case of a semicrystalline thermoplastic/glass fiber (GF) composite, the fiber surface may be sized with the aim to induce a high nucleation density. The closely packed nuclei enforce a primarily outward growth of the crystallites. This leads to a TC interphase layer between the fiber surface and the semicrystalline bulk. Since early reports of transcrystallization [10,11], much effort has been put into the topic and there is still a controversial discussion, as the existence and structure of a TC layer have significant influence on the mechanical properties of the composite [1,4–6,12–17].

The surface free energy is determined by the type of fiber and the sizing or finish components, especially the coupling agent. The nucleation density can be increased by appropriate sizing components at the surface or by matrix additives for the bulk [18–20]. The roughness of the fibers can play a role [5,14,18], especially if epitaxial ordering at the surface is possible [21,22]. The sample temperature influences the nucleation rate at the interface and in the bulk, and the crystallization kinetics [23,24]. Due to the memory effect of incompletely melted polymer crystals, the "temperature history" since the last complete melting plays a role.

Occasionally, the term "transcrystallization" is used in the literature in connection with another phenomenon [25,26]. By applying a shear stress along the fiber axis, a columnar superstructure is formed even atT>Tcr,trans which resembles a TC layer. It was even supposed that TC layers only occur as a result of shearing. However, Wu et al. [27] observed that shear leads to a cylindric crystallization. Varga and Karger-Kocsis [28] determined that this structure does not lead to increased adhesion strength, as the cylindric crystallites are nucleated homogeneously, unlike the heterogeneous nucleation of the TC layer, so no physical coupling occurs.

#### *1.1. Structure of the TC Interphase*

According to Pompe und Mäder [20], three different sections of the TC layer can be specified (Figure 1), which are influenced by the above mentioned parameters.

**Figure 1.** Scheme of the interphase between the fiber and a semicrystalline matrix forming a TC-layer.

Section 1, the nucleation region, includes the interface between fiber and the TC part of the matrix. It is here the nucleation takes place. A TC layer is only formed if the nucleation rate at the surface is higher than in the bulk, e.g., if the crystallization temperature of the interface induced nucleation is higher than that of the bulk crystallization, Tcr,bulk < Tcr,TC. The thickness of Section 1 is determined by the roughness of the interface. The crystal morphology within a layer of the thickness of around the distance of two nuclei is less well ordered compared to the TC layer beyond, where the crystal growth direction is primarily perpendicular to the surface. Within the surface layer, the nuclei grow in a hemisphere, until they are impeded by the neighboring nuclei. For some fiber/matrix combinations, the nuclei themselves have a preferred crystal orientation.

The adjacent Section 2 is the homogeneous part of the TC layer. Its extension is delimited by the impeding bulk spherulites. When its thickness δi,2 is bigger, therefore, the more the TC layer can grow in the time interval between the nucleation at the interface and (a mean point in time of) bulk nucleation. These parameters can be adjusted via a sizing/finish with high nucleating ability, or by nucleating agents/nucleation suppressants within the polymer matrix. The crystal growth rate is determined by the temperature resp. the cooling rate. In the case of slow cooling or isothermal annealing at temperatures T with Tcr,bulk <T<Tcr,trans, an extended TC layer is formed [4–6,8,29]. In samples with high fiber density, the bulk spherulite region can even be fully suppressed [20].

Section 3 (cf. Figure 1) encloses the interface between the TC region and the surrounding bulk, i.e., between the TC and the spherulitic morphology. The thickness δi,3 of this section as well as the specific surface area of the interface are determined by the size of the bordering spherulites [12].

#### *1.2. The TC Interphase and Its Impact on the Micromechanics and Composite Properties*

The influence of a TC interphase on the mechanical properties of the composite is actively and controversially discussed in the literature [26,28–32]. Clark et al. [16] report on polyamide (PA)/GF or PA/carbon fiber (CF) composites with higher strength but lower toughness in the case of transcrystallinity. In another experiment [29,33], PA/GF-composites were cooled at different rates, resulting in TC layers of different thicknesses. The slowly cooled samples with thicker TC layers had a Young's modulus increased by 30% and a bending modulus increased by 70%. However, slowing the cooling rate not only increased the TC layer thickness, but also the degree of crystallization of the matrix and the fraction of α-crystallized PA, making a comparison more difficult.

An intermediate step to study the interplay between the local crystal morphology and the composite properties is the micromechanical study of the influence of transcrystallization on the adhesion strength. The results so far are controversial. Bessel et al. [11] determined a reduced adhesion strength for TC samples. In a fragmentation test, Folkes and Wong [34] determined an increase in the critical fiber length, i.e., also a decrease of the adhesion strength due to TC layers in PP/GF composites. Similar results have been reported by Rolel et al. [35] for polyethylene (PE) matrix composites. In contrast, Carvalho and Bretas [36], Huson and McGill [31] as well as Feldman et al. [14] observed an increase in the adhesion strength due to transcrystallinity, for a number of different fibers in different thermoplastic matrices such as polypropylene (PP), polyethylene (PE) und polyamide (PA). Similarly, Clark et al. [15,16] found high interfacial shear strength and cohesive failure in TC samples, in contrast to adhesive failure of the interface in absence of a TC layer. Hsiao and Chen [30] observed that transcrystallinity had no significant effect on the adhesion strength for a number of composites. The situation is obviously complex, due to interaction and superposition of a number of parameters, influencing the interphase and the adhesion strength. The following paragraphs will discuss these parameters according to the scheme in Figure 1.

Section 1: this section is determined by the fiber interface, i.e., generally by the sizing or finish of the fiber. Sizings (on GF) or finishes (on CF) are used to increase the interfacial strength. Considering transcrystallinity, it is of importance if the surface can induce nuclei for heterogeneous crystallization. The resulting TC layers should act as a "physical coupling" between fiber and matrix [15], increasing the shear strength of this first section. This has been experimentally confirmed [37]. However, usually

the TC is reached via a sizing, which also influences the adhesion. In the present work, GF composites are compared whose sizings are identical except for the film former, with the aim to modify the nucleating properties at comparable adhesion strengths.

Section 2: The homogeneous region of the TC layer influences the composite properties via the properties of the TC morphology as compared to that of the bulk. The degree of crystallization is often higher than in the spherulitically crystallized regions [38]. Karger-Kocsis proposes that in an extended Section 2 the crystallization shrinkage might lead to a weak interface [39]. In the TC layer, the crystallites are smaller and radially aligned in reference to the fiber direction.

There are few experimental results focusing on this region in fiber-reinforced systems. By using specific scanning probe microscopic techniques (phase imaging, nanoindentation), the extension and the Young's modulus of the interphase between GF and PP matrix could be determined [40]. Within a TC layer, the local modulus is higher than in the bulk. Folkes and Hardwick [41] found an increase of storage modulus E' by a factor of 2 and higher shear and tensile strengths for a TC layer compared to a finely spherulitic layer. Marom et al. [38] confirmed these results by dynamic mechanical studies of microtomed sections containing predominantly TC polymer. The TC layer forming samples were more brittle and the fracture energy was smaller than comparable bulk samples. Microbeam synchrotron measurements of TC PP showed that under small load, no change in the microstructure is visible [42]. The authors attribute this to an "anchoring" of the TC layer lamella, confining the strain mainly to the amorphous phase.

Section 3: At the contact of the two crystallization fronts, a second interface is formed in Section 3. This interface is mechanically weak, due to the enrichment of low molecular weight species [43,44] as well as due to low entanglement density [45].

Transcrystallinity occurs in many technically relevant thermoplastics. It has a great influence on composite properties. The existence of a TC layer depends on a number of parameters, on the combination of the materials used as well as on processing conditions. More often than not, the resulting effect of TC morphology is an improvement, but this is by no means unequivocal. The aim of the present work is to study the effect of adhesion strength in PP/GF microcomposites where the fiber sizings are as similar as possible while either inducing a TC crystallization or suppressing it.

#### **2. Experimental**

#### *2.1. Methods*

In the last decade, several pieces of micromechanical equipment have been developed and set up at the IPF to investigate the interphase characteristics between fiber and matrix in composites. The micromechanical experiments include quasi-static pull-out tests, a hysteresis/micro-fatigue test as well as a dynamic test with sinusoidal loads up to 350 Hz. Figure 2 presents a scheme of the experiments. A vessel with matrix and an end-embedded fiber is clamped on an actuator that generates the displacement. The fiber end protruding out of the matrix is glued onto a mandrel that is fixed to a force sensor.

**Figure 2.** General scheme of used micromechanic tests: quasi-static pull out test and hysteresis tests.

The investigation of the adhesion strength parameters between fiber and polymer matrix is carried out with the pull-out setup [46]. Force-displacement curves are obtained by quasi-statically pulling a single fiber out of a polymer matrix. One distinctive feature of the device is the extremely low pull-out speed down to 10 nm/s. Forces between 1 mN to 5 N can be detected by a load cell. The experiment as well as data acquisition, analysis and statistical evaluation are performed by means of a "traditional approach" [47,48].

For the hysteresis and long-term loading test, the fiber is periodically loaded and the hysteresis loop is analyzed for the failure behavior [49]. The general setup (cf. Figure 2) is optimized for hysteresis measurements (low frequency range: 0–10 Hz) of single-fiber model composites, especially for long-term tests. A piezotranslator in combination with a piezoresistive load cell provides a zero backlash deflection and a long-term stable force measurement. Measurement data are the amplitude of the measured value, hysteresis, root point drift as well as elastic and loss energy. In long-term load tests (e.g., Wöhler, relaxation or load increase tests), the periodic change of the phase angle between force and elongation and the periodic change of stiffness and damping are analyzed.

An atomic force microscope AFM (Bruker, dimension) was used in the tapping mode to characterize the fiber surface and fracture surfaces.

Polarization microscopy is used to study the crystallization behavior of the PP matrix. Matrix material and separated single fibers are sandwiched between two glass plates. The thickness of the layer is determined by spacers (150 μm). The sandwiches are observed by polarization microscopy (Scope A1, Carl Zeiss, Oberkochen, Germany) using a hot plate (LTS350, Linkam, Tadworth, UK). The sandwich structures were heated to 192 ◦C for five minutes to fully melt the matrix. Then they are cooled to the crystallization temperature at 60 K/min.

#### *2.2. Materials*

GF were spun at the IPF spinning device. Yarns of 204 filaments were spun and sized with aqueous sizings. Two sizings are compared:

On the one hand APS-PP containing gamma aminopropyl triethoxisilane (APS) as coupling agent, and a PP film former (maleic acid anhydride grafted PP, MaPP), and on the other hand APS-PU with the coupling agent APS and a polyurethane (PU) film former. They are referred to as APS-PP and APS-PU fibers, respectively. As matrix PP, HD 120MO PP homopolymer (Borealis, MFR 8 g/10 min), a typical injection molding grade, and HH 450 PP homopolymer (Borealis MFR 37 g/10 min) a typical grade for fiber spinning applications were used, and are referred to as hiMW and lowMW PP. The two matrices were functionalized by either 0.5 or 2.1 wt % MaPP Exxelor PO 1020 (ExxonMobil Corp., Irving, TX, USA).

Fibers and matrix were compounded on a twin screw extruder and injection molded into standardized dog bone shaped specimens. The average fiber length was approx. 500 μm in all samples.

Tensile tests were conducted on the Universal Testing Machine Zwick 1456 (Zwick-Roell, Ulm, Germany), according to ISO 527-2/1A/50. Unnotched Izod impact tests were conducted with HIT 50P, Zwick/Roell, Germany according to ISO 179/1eU at room temperature and 50% relative humidity. Each value obtained represents the average of ten specimens.

The single-fiber model composites for quasi-static pull-out, hysteresis and cyclic loading investigations are all prepared in the same manner by using the IPF-made fiber embedding device [50] An 80 mg drop of matrix is melted. The pneumatically held fiber is positioned above the matrix with the aid of two long-distance microscope optics, monitoring the xy-positioning resp. the z-coordinate of the fiber above the matrix. The fiber is embedded with a micro drive (embedding length: 20–1000 μm with 0.1 μm accuracy—however, due to the formation of a meniscus, the embedding length may vary and is therefore optically determined after each pull out test). The complete temperature pattern is controlled by a PC (room temperature (RT) to 400 ◦C with 1 K accuracy). The chamber may be flushed with Argon.

#### **3. Results**

Figure 3 shows polarization micrographs of the single fiber in a neat HH450 PP film. Fibers were heated (10 K/min) to 210 ◦C for 0.50 min, then cooled at 200 K/min to the isothermal temperature Tiso = 135 ◦C, at which the crystallization behavior was observed.

**Figure 3.** Polarization micrographs of single glass fibers in PP, Tiso = 135 ◦C, scalebar 100 μm. Top image: APS-PU sized fiber, bottom image: APS-PP sized fiber in HH450 PP matrix. The APS-PP sized fiber induces a TC layer.

Between crossed polarizers, the spherulites in the bulk PP film are clearly visible. Around the APS-PP sized fiber, a TC film of about 3 μm thickness is formed. This confirms an earlier differential scanning calorimetry study of the two sized fibers in PP, which showed that the APS-PU sized fiber induces no TC layer, whereas the APS-PP sized fiber induces a distinct TC layer [20].

The size of the spherulites for the two neat polymers is comparable. Adding 2% MaPP increases the nucleation density and therefore reduces the spherulite size, the crystallization is speeded up: the half time t50% is reduced by 15%.

A series of GF/PP composites were compounded and injection molded with varying MaPP content, with fiber sizing known to induce or suppress a TC layer and differing PP molecular weight. Table 1 shows the mechanical properties of injection molded GF/PP specimens. In GF/PP composites, usually a small percentage of MaPP is added to improve adhesion as the maleic moieties form covalent bonds with the APS or weaker bonds with the OH on the glass surface. A side effect of MaPP is the nucleation of spherulites in the bulk. This side effect reduces the thickness of the TC layer. Therefore, the mechanical properties are determined at two concentrations of MaPP, a technically relevant 2% and a low 0.5%. For both PP grades, there is a slight improvement of tensile strength due to the TC in the 2% MaPP samples. In the case of 0.5% MaPP specimens, the effect of the TC layer is far more pronounced (albeit at a lower mechanical level). The toughness is doubled in the case of the TC interphase in high MW PP. The Young's modulus is increased strongly upon addition of GF, but does not depend on the TC interphase.

**Table 1.** Mechanical parameters of injection-molded specimens of GF reinforced PP.

Table 2 shows the results of micromechanical tests on single-fiber model composites. For both matrix PP grades, the local shear strength *τd*, the critical energy release rate of the interface, *Gic* and the frictional shear stress after debonding, *τf*, are greater for the TC composites. In this interphase sensitive method, the effect is clearly seen even in composites with 2% MaPP. AFM images of the pulled out fibers show clear differences in the fiber fracture surface, cf. Figure 4. For a TC interphase, the fracture surface is rough, with 500 nm structures dominating the surface morphology. In absence of a TC layer, no such structures are seen. These structures are interpreted as either the outer surface of the TC layer (Section 3 in Figure 1) or the sized fiber surface. A model for the fracture is proposed (right hand sind in Figure 4) where in the case of a TC interphase, the failure occurs outside the TC layer whereas without TC layer it occurs at the fiber surface.

**Table 2.** Micromechanical pull out tests results on GF /PP samples with 2.0% MaPP, parameters local shear strength, *τd*, the critical energy release rate of the interface, *Gic*, and the frictional shear stress after debonding, *τf .*

**Figure 4.** AFM height images of the fracture surfaces of single-fiber model composites after pull-out with HH450, Top: fibers inducing TC interphase, bottom: fibers suppressing a TC interphase. The height and size of the protrusions is 400 nm × 30 nm for the TC fracture surface and 50 nm × 8 nm for the non-TC fracture surface. A model for the fracture (right) is proposed where in the case of a TC interphase, the failure occurs outside the TC layer whereas without TC layer it occurs at the fiber surface.

In addition, micromechanical dynamical [51] and micromechanical hysteresis measurements [49] were performed on single-fiber model composites at two different embedded lengths, (200 μm, 600 μm).

Figure 5 shows the results of the hysteresis measurements for the four fiber matrix combinations high vs. low MW and TC vs. non-TC interphase for 600 μm embedded length. All samples remain intact with only some degradation. As the embedded length and fiber diameter vary, the value of the force as well as the stiffness (force/displacement) is subjected to error. However, the variation of stiffness with cycle number provides information on the degradation of the interphase, as well as the width of the hysteresis curve or the area included in the hysteresis loop, which is equal to the inelastic energy loss during the cycle (Figure 6). The stiffness of the high MW interphases is reduced by 20%, whereas the stiffness of the low MW samples is reduced by 40% at 40,000 cycles. This is independent on the interphase crystallization. The width of the hysteresis loop at zero force is initially higher for high MW than for lower MW, and for non-TC interphase samples than for the TC ones. Energy loss processes are higher for high MW and in the non-TC interphase. With cycle number, the non-TC samples hysteresis decreases whereas the TC interphases have a constant hysteresis width. Some interphase deterioration processes seem to occur in non-TC samples that are suppressed in samples with a TC interphase.

**Figure 5.** Hysteresis loops of single-fiber model composites loaded with amplitudes of ±1.8 μm at a frequency of 10 Hz (sized glass fibers APS/PP- or APS-PU-embedded in either PP HD120M or PP HH450) after the given number of cycles.

**Figure 6.** Data evaluated from the hysteresis curves: Relative stiffness (**left**) and width (**right**) of the hysteresis loops. Lines are guide to the eye.

If a smaller embedding length of 200 μm is chosen, the non-TC samples fail, whereas the transcrystalline samples remain intact up to 400,000 cycles.

The AFM images of either APS-PP or APS-PU fibers before embedding show a relatively smooth, homogeneous surface (Protrusions height <30 nm, diameter <100 nm, phase difference <15◦). When the fiber is pulled out of the matrix after cyclic loading, the fracture surfaces are significantly more inhomogeneous (Figure 7). On the APS-PP surfaces, structures of the order of magnitude of 500 nm are found that are strongly oriented along the fiber (=load) axis. This longitudinal alignment is also found in the phase image. In the APS-PU sized fiber, the structures are significantly larger (2 μm), and the contrast in the phase image is much higher.

**Figure 7.** Tapping mode AFM images of fiber surfaces fractured after cyclic tests on single-fiber model composites.

For the APS-PP sized fibers, a TC structure is expected, leading to a fine-grained structure on the surface. The crystals are strongly deformed under load, but there is no material contrast, so the fracture surface is within the PP. In the APS-PU, an alignment of the crystallites is also observed, mainly in the phase image. The phase image also indicates a surface with varying material parameters, i.e., a fracture surface that is located partially in the PP/PU interphase and partially within the PP matrix: a number of larger spherulites adhere to the surface. They might be the cause for the higher friction seen in the micromechanical measurements.

#### **4. Conclusions**

A transcrystalline interphase has in the past been observed to improve ]mechanical properties, but contradictory results have also been observed. In the present study, an enhancing effect of the transcrystalline layer is shown. However, it may be reduced due to the effect of additives such as MaPP, which is added as a coupling agent between fiber and PP matrix. As MaPP will also induce nucleation in the bulk, it reduces the effect of the transcrystalline layer. For standard MaPP concentration (2%), the effect of the remaining TC layer is weak, tensile strength and toughness are only increased by 5 to 10%. If only 0.5% of MaPP are added, the TC layer is more pronounced, leading to a 70–100% increase of strength and a 25 or 125% increase in toughness, depending on the MW of the matrix polymer. In single-fiber model composite interface-specific tests, such as the single-fiber pull out test or single fiber hysteresis test, the interphase enhancement due to the transcrystallization is evident even at 2% MaPP. For both matrix PP grades, the local shear strength, *τd*, the critical energy release rate of the interface, *Gic*, and the frictional shear stress after debonding, *τf*, are greater for the TC composites. AFM images reveal a failure at the sizing layer for non-TC samples and failure at the outside of the TC layer in TC model composites.

**Acknowledgments:** The authors are indebted to Steffi Pressler and Alma Rothe for technical assistance, and to Wolfgang Jenschke for helpful discussions and for improving the micromechanical test setups. The authors thank the German Research Foundation DFG for financial support (Project MA 2311/1-1).

**Author Contributions:** Hanna Brodowsky and Edith Mäder studied the literature, conceived and designed the experiments. Hanna Brodowsky performed the experiments and analyzed the data; Hanna Brodowsky and Edith Mäder discussed the results, Hanna Brodowsky wrote the paper.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

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