Mechanical and interfacial properties of phenolic composites reinforced with treated cellulose fibers.
Phenolic resins are used in a wide variety of applications such as adhesives, surface coatings, powders of molding, thermal insulations materials, and composites. Phenolic matrix composites largely inherit the characteristics of phenolic resins, and thus the mentioned composites have good resistance to high temperatures, fire, shock, abrasion and chemical attack, excellent electrical characteristics, and dimensional stability , Low mechanical performance is the main drawback of these phenolic materials. Several efforts have been made in order to improve mechanical properties of composites, which traditionally consisted of incorporating glass or aramid fibers as reinforcement. In recent years, concern for the environment has led researchers to substitute synthetic fibers by natural ones [2, 3].
Cellulose fibers are a potential raw material for making environmental friendly and low cost composites. Cellulose fibers are more resistant to the absorption of noise, lighter, and lower cost than conventional glass fibers. In addition, cellulose fibers can be obtained and recycled using less energy. Apart from naturally occurring plant fibers, cellulose fibers can be industrially obtained in form of man-made or regenerated fibers using the viscose or the NMMO (N-methylmorpholine-A-oxide) process. The potential of regenerated fibers as reinforcement agents in composites is being recognized, due to their high purities, few defects, and good reproducibility of properties , Some works with regenerated cellulose fiber-reinforced composites have been recently reported, although most of them with thermoplastic matrices, as polylactic acid (PLA), polyoxymethylene (POM), poly(3-hydroxybutyrate) (PHB), or polypropylene [5-8], In this study, regenerated viscose fibers are incorporated as reinforcement into thermoset phenolic composites.
The major limitations of using cellulose fibers as reinforcement in composites are their poor resistances to moisture absorption and bad adhesions with most of matrices. Surface modification of the fibers is necessary to solve the adhesion problems and obtain better mechanical properties in these reinforced materials. Several fibers' treatments have been reviewed in the literature, including reaction with acid compounds and anhydrides, alkali treatment, coupling with organosilanes, chemical grafting, and surface activation by physical agents [9-12], The treatment with a sodium hydroxide solution is one of the most widely used methods. It removes impurities of the fibers, producing a rough surface and improving the fibers adhesive characteristics [11, 13]. Moreover, this treatment increases the available surface area of the fibers for the join with the matrix. The employment of silanes as coupling agents between fibers and matrix stands out as one of the most efficient chemical methods for treating fibers [14-17]. This treatment improves the fibers-matrix interfacial bonding and makes the fibers more hydrophobic, limiting the moisture absorption of the reinforced composites. The effects of alkali-silane combined treatments have begun to be studied in epoxy and polyester composites, but the mechanisms are not well-known .
The aim of this study is to establish the most appropriate preparation of cellulose fiber-reinforced phenolic composites according to the modification of the fibers incorporated to these materials. Fibers treatments with sodium hydroxide and/or silanes are studied. Previously, it is necessary to optimize the curing process of the phenolic matrix in order to apply the optimal operating conditions to the curing of composites. In both studies, curing of phenolic matrix and preparation of composite, the results are referred to the mechanical and morphological properties of the obtained materials.
MATERIALS AND METHODS
A thermoset phenolic-based water resol resin and an organic ester hardener (ACE-1035) used in the matrix preparation of the composites were supplied by Momentive Specialty Chemicals. Viscose cellulosic fibers from eucalyptus wood of length 1.7-38 mm and linear density 1.7 dtex (0.17 g/1000 m) used as reinforcement were supplied by Sniace.
Analytical grade sodium hydroxide (NaOH > 98%) was employed in the alkali treatment of the fibers (Panreac Quimica) and two organosilanes, namely (3-aminopropyl) trimethoxysilane (APS) and 3-(2-aminoethylamino) propyltrimethoxysilane (AAPS), were employed as coupling agents (Sigma-Aldrich[R]).
Experimental Set-Up and Procedure
Figure 1 shows the experimental scheme followed in this work. Polymer matrix was prepared by mixing the phenolic resol resin and the ester hardener in the ratio 80:20 by weight with mechanical stirring at 800 rpm for 2 min. Then, the mixture was poured into a steel mold (160 X 110 X 6 mm) and cured in an oven at different temperatures and times. The study of the phenolic matrix curing process was carried out using a factorial experimental design with two factors and two levels, designated by [2.sup.2] (11 runs: [2.sup.2] + 3 central points + 4 star points). The variables studied were curing temperature (T) and time (t), and their operating levels were 50-90[degrees]C and 1-4 h, respectively. The responses measured for model development were the tensile and flexural properties: strength ([sigma]), elongation ([epsilon]), modulus (E), and strain energy density (SED). The experimental conditions and results obtained for the curing of phenolic matrix are shown in Table 1. Letters "T" and "F" are used to differentiate properties in tensile and flexural tests, respectively. Data processing was accomplished using Statgraphics Centurion XV, which enables one to apply analysis of variance (ANOVA) and multiple linear regression methods. The effect of curing temperature and time on the morphology of the phenolic matrix was also studied. In addition, the sample cured under optimal conditions was dried in a vacuum oven at 70[degrees]C up to reach a constant weight. The change on the mechanical properties of the material after the drying stage was further evaluated. The drying of phenolic material is frequently recommended due to its common high water content [19-21].
The preparation of the cellulose fiber-reinforced phenolic composites was carried out by mixing first the resol phenolic resin and the cellulose fibers (800 rpm, 15 min), and then the hardener (800 rpm, 2 min). All composites were reinforced with 3 wt % of cellulose fiber, cured at the optimal conditions selected in the curing study of the phenolic matrix, and dried in a vacuum oven at 70[degrees]C up to reach a constant weight. Composites were obtained with untreated fiber and with fibers modified by alkali treatments, silane treatments, and combined treatments alkalisilane. To select the composite more appropriate, mechanical and morphological properties were studied. Table 2 summarizes the treatments applied to the fibers in this study and the code names of the resulting randomly oriented cellulose fiber-reinforced phenolic composites. Conditions for each one of the fibers treatments were:
* Alkali treatments. The fibers were soaked in 1 or 5 wt % NaOH solution for 2 h at room temperature. They were further washed until pH = 7, and finally dried at 60[degrees]C for 24 h.
* Silane treatments. The cellulose fibers were treated for 120 min with 2.2% of APS silane or 100 min with 1.5% of AAPS silane, using an acidified 80/20 wt/wt ethanol/ water medium for the silanes hydrolysis. Finally, the fibers were dried at 60[degrees]C for 24 h.
* Combined treatments alkali-silane. Cellulose fibers were treated with NaOH and then with one of the silanes in the conditions detailed above. Four combined treatments were applied, which result of different combinations of alkali and silanes treatments.
Mechanical and Morphological Characterization
Mechanical properties of the phenolic matrices and cellulose fiber-reinforced composites were evaluated. Test samples were cut using a computer numerical controlled or CNC milling machine (Galdabini ICP 3020) and then polished with a Buehler MetaServ[R] 3000 polisher to a thickness of 3.2 mm. Tensile and flexural tests were performed using a universal testing machine (Zwick/Roell Z030) in accordance with ASTM standards D 638 and D 790, respectively. Tensile and flexural strength, elongation, modulus, and strain energy density were obtained from the respective stress-strain curves, using at least five specimens per test condition.
A scanning electron microscope (SEM, JEOL JM-6400) was utilized to observe the morphology of the fracture surfaces of the matrices and the cellulose fiber-phenolic matrix interfaces of the composites. Square bars with a 10 [mm.sup.2] cross-section and 40 mm in length were cryogenically frozen to avoid structural deformations of the material and then were fractured. Gold sputtering onto the fractured surfaces of the samples was used to impart them electrical conductivity. The operation voltage of the SEM was 40 kV.
Extent of dispersion of the fibers into the matrix for the different composites was analyzed with an X-ray diffraction instrument (Philips X'Pert-MPD). Cu K[alpha] radiation with K[[alpha].sub.1] = 1.54056 [angstrom] and K[[alpha].sub.2] = 1.54439 [angstrom] was used. Powder cellulose fiber, phenolic matrix, and composites were scanned over the interval of 2[theta] from 2 to 40[degrees] with a step size of 0.02[degrees] and a time per step of 2 s. The accelerating voltage was 45 kV. Using X-ray diffraction (XRD), the effect of the fiber treatment on the intercalation behavior of these fibers in the composites was analyzed. The fiber-fiber interlayer spacing (d) was calculated using the Bragg equation: [lambda] = 2d sin [theta], where [lambda] is the wavelength.
RESULTS AND DISCUSSION
The measured values of strength, elongation, modulus, and strain energy density provided a basis for evaluating the behavior of the cured phenolic matrix. The values of these tensile and flexural responses for the different combinations of curing temperature and time of the phenolic matrix are summarized in Table 1.
The models for maximum strength obtained after performing the analysis of variance (ANOVA) and excluding the non-significant effects for a significance level of 95% (P > 0.05% and/or F < 18.51) in tensile and flexural tests are given by Eqs. 1 and 2, respectively. The contour map of the tensile strength as a function of curing temperature and time is shown in Fig. 2a. Plot exhibits a maximum value of 8.4 MPa, which can be obtained with different combinations of curing temperature (67.5-87.5[degrees]C) and time (2-3.5 h). Tensile strength of the matrix decreases at curing temperatures above 87.5[degrees]C and/or times over 3.5 h due to the formation of bubbles in the material , Flexural strength of the cured matrix varies in the range 11.7-13.7 MPa, as shown in Fig. 2b. Flexural strength increases with the curing temperature for any time studied. For a given temperature, flexural strength increases with the curing time up to a value (different for each one of temperatures) and then it remains constant and even reduces.
[MATHEMATICAL EXPRESSION NOT REPRODUCIBLE IN ASCII] (1)
[MATHEMATICAL EXPRESSION NOT REPRODUCIBLE IN ASCII] (2)
The models obtained for elongation of the phenolic matrix in tensile and flexural tests are given by Eqs. 3 and 4. The contour maps for the elongation of the phenolic matrix in tensile and flexural tests are shown in Fig. 2c and d, respectively. Tensile and flexural elongation increased when the curing temperature and/or time were decreased, reaching maximum values for tensile and flexural elongation of 5.3 and 5.8%, respectively. When the curing is performed at low curing temperatures and times, the material retains more plasticity at the expense of not increasing its stiffness . The behavior of the material varies from brittle to ductile with the increment of its elongation before breaking. This also results in the reduction of its strength (Fig. 2a and b).
[[epsilon].sub.T](%) = 8.10394 - 0.0476847 * T - 0.201796 * t [R.sup.2] = 97.85% (3)
[[epsilon].sub.F](%) = 8.03891 - 0.040357 * T - 0.191387 * t [R.sup.2] = 89.08% (4)
The elastic modulus of a material is related to its cross-linking degree and stiffness. The models obtained for this response in tensile and flexural tests after performing ANOVA analysis and excluding non-significant effects are given by Eqs. 5 and 6, respectively. The contour map of the elastic modulus in tension for the cured matrix is shown in Fig. 2e. Increasing the curing temperature and time produces the non-linear increment of the tensile modulus. The range of phenolic matrix curing times between 2.75 and 3.25 h is the most interesting region in this map due to maximum values of the tensile modulus are reached. Thus, for a curing temperature of 65[degrees]C, as example, tensile modulus rises from 255 to 280 MPa, by prolonging curing time from 1 to 3 h. Above 3.25 h, tensile modulus is reduced because of air bubbles are formed in the material , This decrease is consistent with the reduction of the tensile strength over a similar curing time described above. The contour map for the elastic modulus in flexion of the matrix is shown in Fig. 2f. Increasing curing temperature and/or time leads to the increment of the flexural modulus, whose values vary between 200 and 300 MPa in the variables studied ranges.
[E.sup.T](MPa) = 111.21 + 1.70316 * T + 38.1685 * t - 6.16477 * [t.sup.2] [R.sup.2] = 94.86% (5)
[E.sub.F](MPa) = 85.4584 + 2.01102 * T + 10.579 * t [R.sup.2] = 91.53% (6)
The models obtained for strain energy density of the cured matrix in tensile and flexural tests are given by Eqs. 7 and 8. SED represents the integral of the area under the stress-strain curve and provides a measure of the energy absorbed by the material during the test and thus of its resistance to the fracture [24-26]. The contour maps for SED of the matrix in tensile and flexural tests are shown in Fig. 2g and h, respectively. In both cases, SED increased when the curing temperature and/or time were decreased. This behavior is similar to that described above for the elongation. SED in tensile and flexural tests reaches maximum values of 2.5 X [10.sup.5] and 3.8 X [10.sup.5] J/[m.sup.3], respectively.
[MATHEMATICAL EXPRESSION NOT REPRODUCIBLE IN ASCII] (7)
[SED.sub.F](J/[m.sup.3]) = 577,396 - 3,105.05 * T - 20,329.2 * t [R.sup.2] = 91.92% (8)
The selection of the optimal conditions for curing process of the matrix was based on obtaining good mechanical properties of the material in tensile and flexural tests, using curing temperatures and times as mild as possible. After analyzing separately all the contour maps obtained (Fig. 2), optimal curing conditions were selected fundamentally attending to the tensile strength map, which showed a maximum (8.4 MPa) in the studied range of the variables (Fig. 2a). Thus, the optimal curing temperature and time were found to be 75[degrees]C and 2.75 h, respectively, which led to the mentioned maximum in tensile strength with fairly mild conditions. In turn, these conditions represent a compromise solution between high values of strength and elastic modulus of the matrix (which increased when the curing temperature and/or time were increased), and acceptable values of elongation and SED (which decreased when the curing temperature and/or time were increased).
SEM micrographs of fracture surfaces of the phenolic matrix cured in some of the conditions of the experimental design are shown in Fig. 3. Samples were carefully selected in order to confirm the influence of curing temperature (Fig. 3a, c, and d) and time (Fig. 3b, c, and e) on the morphology of the phenolic matrix. Figure 3c shows an intermediate behavior of the matrix in the studied temperature and time ranges. SEM image of the phenolic matrix cured under optimal conditions is also included (Fig. 30- An increment on the number of bubbles occluded in the material is appreciated when high temperatures (98.3[degrees]C) or prolonged times (4.6 h) are applied during the curing (Fig. 3d and e, respectively). These bubbles are the responsible of the tensile strength reductions for high temperatures and/or prolonged times described above, and they have been reported in the literature for similar matrices , Moreover, deformation lines are observed when curing temperature or time is increased, due to fracture in the material changes from "plastic" (Fig. 3a and b) to "fragile type" (Fig. 3c-f) with the cross-linking progress . The morphological images of the phenolic matrices confirm the optimal conditions mentioned above for the curing process of these materials. Not bubbles appear in the image of the sample corresponding to these conditions (Fig. 3f).
The theoretical mechanical properties for the optimal curing conditions can be predicted employing the models represented by the Eqs. 1-8. The phenolic matrix was cured under optimal curing conditions, and predicted and experimental values of the mechanical properties were compared in order to validate the models (Table 3). The difference between the experimental and the predicted values was lower than 10%, confirming the predictive accuracy of the proposed models and validating the predictions. Note that strengths in tensile and flexural tests of the matrix after the incorporation of the final drying stage were 156 and 189% of the strengths of the undried matrix, and elastic moduli were three-fold (Table 3). This drying stage enabled us to obtain a competitive phenolic matrix in comparison with other literature materials, despite of working with a high water content based resin. For example, while strength values in tensile and flexural tests of the phenolic matrix in this study were 13.28 and 24.20 MPa, Joseph et al.  and Sreekala et al.  found tensile strength values of 7 and 10 MPa, respectively, and flexural strength values of 10 MPa, also for phenolic matrices.
Cellulose Fiber-Reinforced Phenolic Composites
The results of tensile and flexural tests for the studied unreinforced and reinforced composites are shown in Fig. 4. Dot and dash horizontal lines on the figure are the mechanical properties in tension and flexion, respectively, of the unreinforced reference material (phenolic matrix). Note that the coefficient of variation (ratio of the standard deviation to the mean) is lower than 10% in all cases. Untreated fiber-reinforced composites (Cell-Ph) present better mechanical properties in tension than the reference matrix (Unreinforced), due to cellulose fibers support an important part of the material stress and increase its stiffness . However, mechanical properties in flexion are reduced with the addition of cellulose fibers into the phenolic matrix, which is a result of the formation of voids in the material [27, 30]. The appearance of these voids will be proved later with the SEM images. The incorporation of cellulose fibers in the matrix reduces the flexural properties due to the presence of voids in the material has a greater effect on these properties than on the tensile ones .
Alkali treatment of the fibers with low concentration of NaOH (1%) improves the interfacial bonding fiber-matrix and provides additional mechanical attachment sites, enhancing some of the mechanical properties of the composites (Cell-1%NaOH-Ph) [31-34], It can be calculated from data of Table A1 that tensile modulus and flexural strength, modulus, and SED improve 4, 12, 45, and 9%, respectively, in relation to the reference sample properties. In contrast, alkali treatment of the fibers with 5% NaOH reduces the mechanical properties of the composites (Cell-5%NaOH-Ph). The cellulose fibers can weaken and even degrade over a concentration of NaOH, and thereby reduce the mechanical properties of the resulting composites , Degradation of the viscose fibers with the alkali treatment has been described in a previous work for NaOH concentrations over 10% , and weakening of the fibers could begin even at lower concentration. This will be proved later with the SEM images.
Treatment of the cellulose fibers with APS and AAPS silanes enhances also the tensile properties of the composites, but advances were not found with the alkali-silane combined treatments. The most significant improvements obtained from data of Table A1 were 30 and 63% for the tensile elongation and SED with the APS silane treatment (Cell-APS-Ph), and 25, 52, and 110% for tensile strength, elongation, and SED with the AAPS silane treatment (Cell-AAPS-Ph). Regarding the flexural properties of the composites, progresses were only observed with the AAPS silane treatment. Thus, the fiber treatment with AAPS silane led to the best mechanical properties of the composites.
SEM micrographs of the cellulose fibers-phenolic matrix interfaces of the composites, before and after treatment of the fibers, are shown in Fig. 5. Images of the raw fiber-reinforced composites, Cell-Ph, at different magnifications (Fig. 5a and b) reveal voids between the cellulose fibers and the phenolic matrix, evidencing the poor adhesion between the untreated viscose fibers and the matrix. Weak or no adhesion is more frequently found in smooth surface regenerated fibers than in natural ones, whose rough surfaces improve the adhesion to the matrix , Furthermore, debonding appears in the Cell-Ph composite of the Fig. 5a, which is the result of a poor stress transfer between matrix and fibers .
In the image of the 1% NaOH treated fiber-matrix interface (Fig. 5c), a reduction on the interfacial void size is observed, compared with the image of the untreated fiber (Fig. 5b). Alkali treatment at this concentration increases surface roughness of the fibers, resulting in better fiber-matrix interfacial adhesion in the composites . This enhanced interfacial adhesion is responsible for the reduction in the interfacial void size and the improvement of some of the mechanical properties described above. Liu et al.  described a similar behavior after treating jute fibers with 2% NaOH and incorporating them as reinforcement into poly(butylene succinate) matrix composites.
Larger voids can be appreciated in the image of the 5% NaOH treated fiber-matrix interface (Fig. 5d). The fiber breakage observed in this image is characteristic of a damaged or a weakened fiber. Moreover, other evidences of fiber degradation caused by the solution of 5% NaOH can be observed in the detail of the figure. As we mentioned earlier, fibers can be weakened at high alkali concentrations, resulting in reductions in interfacial adhesion and mechanical properties of derived composites . Voids around the cellulose fibers derived of a poor interfacial adhesion act as stress concentration points and are the responsible of the reduction of mechanical properties of the composites ,
Interfaces of the APS and AAPS silanes treated fibers-reinforced composites (Cell-APS-Ph and Cell-AAPS-Ph) are shown in Fig. 5e and f, respectively. Interfacial voids and debonding completely disappear in both cases, proving the improvement in the fiber-matrix adhesion caused by the silane treatments. Furthermore, it can be appreciated that part of the phenolic matrix remains adhered to the surface of the cellulose fibers after the rupture, especially with the AAPS silane treatment, proving the formation of cohesive interfaces. Other authors have observed by SEM improvements in the interfacial properties of several fibers-matrix systems when silanes are employed as coupling agents [40-42].
Regarding the influence of the combined treatments alkali-silane of the fibers on the interfacial properties of the composites, two different behaviors have been found. On the one hand, the fiber treatment with 1% NaOH and any of the two silanes improves the compatibility fiber-matrix (Fig. 5g and h). This behavior matches with the improvements in fiber-matrix adhesion previously discussed when the treatments with 1% NaOH, APS, or AAPS silanes were applied separately. However, the apparent good interfacial adhesion is not consistent with the poor mechanical properties of the respective composites shown in Fig. 4 (Cell- 1% NaOH + APS-Ph and Cell-1% NaOH + AAPS-Ph). Although the mentioned combined treatments improve the fiber-matrix adhesion (SEM observations), they may damage the fibers and thereby reduce the mechanical properties of the resulting composites. Megiatto et al.  described a similar disparity between morphological and mechanical properties in composites. On the other hand, the fiber treatment with 5% NaOH and APS or AAPS silanes leads to a poor interfacial adhesion fiber-matrix (Fig. 5i and j). Large voids in the interface can be observed, similar to that of 5% NaOH treated fiber-composites. These voids are the responsible for the poor mechanical properties of the composites Cell-5% NaOH + APS-Ph and Cell-5% NaOH + AAPS-Ph described above.
Figure 6 shows the X-ray diffractograms of the cellulose fibers, the phenolic matrix (Unreinforced), the untreated fiber-reinforced composite (Cell-Ph), and the composites reinforced with fibers modified by some of the methods studied in this work. The combine treatment of the fibers with 1% NaOH and APS silane (Cell-1% NaOH + APS-Ph) has been selected to represent the behavior of all the combine treatments. The diffractogram of the viscose cellulosic fiber showed the typical X-ray diffraction pattern of regenerated cellulose (Cellulose II) with major peaks at around 2[theta] of 12.7[degrees] (101), 20.9[degrees] (10[bar.1]), and 22.1[degrees] (002) . The peak appearing at 12.7[degrees] (101) was selected to follow the extent of dispersion of the fibers into the matrix and it did not overlap with any peak of the phenolic matrix. The peak was shifted to a lower angle when the fibers were intercalated with the matrix in the composites.
The interlayer spacing (d) between fibers was calculated using the Bragg equation, to verify a homogeneous dispersion of the fibers in the polymeric matrix (so-called intercalation). In addition, the relative intercalation (RI) of the polymer in the cellulose fibers was calculated as the difference between the d-spacing of the cellulose in the composites and the spacing of the cellulose layers in the pure sample, relative to the latter. The d-spacing and relative intercalations of the cellulose fibers in the composites are listed in Table 4. The extent of intercalation of samples depends on the treatment of the fibers which are incorporated into the composites. The peak of the (101) lattice plane was shifted to a lower angle with the treatments of the fibers with the silanes APS and AAPS. This shift corresponded to a higher d-spacing between the fibers in the composites and to a higher relative intercalation. Silanes APS and AAPS improved the compatibility between fibers and matrix, which increased the distance among two different fiber layers (d-spacing). In other words, silane treatment of the fibers, especially AAPS treatment, improved the dispersion and intercalation of the cellulose fibers in the phenolic matrix. The rest of the treatments did not cause significant changes in the dispersion of the fibers in the polymeric matrix.
In short, the AAPS silane treated cellulose fiber-reinforced phenolic composite was found to be the material with the best mechanical and morphological properties. AAPS silane acts as an effective coupling agent between the viscose cellulosic fibers and the phenolic matrix. It improves the dispersion and intercalation of the fibers in the matrix, enhances their compatibilities and adhesions, the formation of cohesive interfaces, and the mechanical properties of the composites.
The optimal curing conditions established for the phenolic matrix were 75[degrees]C and 2.75 h, which represent a compromise solution between obtaining high values of strength and modulus and acceptable values of elongation and SED of the material. The statistical models obtained were suitable for reproducing the experimental data of the matrix curing. The incorporation of a final drying stage in the phenolic matrix preparation of 70[degrees]C and vacuum provided a significant improvement on the bulk mechanical properties of this material.
After the mechanical and morphological characterization of the cellulose fiber-reinforced phenolic matrix composites, it can be concluded that the treatment of the fibers with 1.5% of 3-(2-aminoethylamino) propyltrimethoxysilane (AAPS) for 100 min provided the reinforced composites with the best properties. The AAPS silane acted as an effective coupling agent between the viscose cellulosic fibers and the phenolic matrix, enhancing the mechanical properties of the final composites in tension and flexion. Furthermore, this treatment improved the fiber-matrix compatibility, led to the formation of cohesive interfaces between both materials, and improved the dispersion of the viscose cellulosic fibers in the phenolic matrix.
TABLE A1. Mechanical properties values in tensile and flexural tests of the unreinforced and 3 wt % cellulose fiber-reinforced phenolic composites. [sigma] (MPa) [epsilon] (%) Composites T F T F Unreinforced 13.28 24.20 1.86 2.31 Cell-Ph 15.09 23.39 2.35 2.17 Cell-1% NaOH-Ph 13.31 26.99 1.79 2.13 Cell-5% NaOH-Ph 7.53 19.12 1.27 2.08 Cell-APS-Ph 14.71 23.43 2.41 2.15 Cell-AAPS-Ph 16.65 25.13 2.82 2.39 Cell-1% NaOH + APS-Ph 10.30 21.15 1.46 2.20 Cell-1% NaOH + AAPS-Ph 10.96 17.34 1.65 2.03 Cell-5% NaOH + APS-Ph 11.97 20.96 1.89 2.14 Cell-5% NaOH + AAPS-Ph 11.66 20.08 1.96 2.37 E (MPa) SED (J x [10.sup.-5]/ [m.sup.3] Composites T F T F Unreinforced 835.04 689.46 1.32 2.53 Cell-Ph 887.72 654.26 2.10 2.25 Cell-1% NaOH-Ph 868.11 1000.39 1.28 2.75 Cell-5% NaOH-Ph 660.86 656.57 0.51 1.77 Cell-APS-Ph 872.04 596.32 2.15 2.20 Cell-AAPS-Ph 870.48 726.86 2.77 2.73 Cell-1% NaOH + APS-Ph 829.05 621.90 0.82 2.08 Cell-1% NaOH + AAPS-Ph 773.55 682.61 0.96 1.60 Cell-5% NaOH + APS-Ph 752.19 630.00 1.21 2.02 Cell-5% NaOH + AAPS-Ph 713.46 681.09 1.24 2.21
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Ester Rojo, Mercedes Oliet, M. Virginia Alonso, Belen Del Saz-Orozco, Francisco Rodriguez
Department of Chemical Engineering, Complutense University of Madrid, Avda. Complutense s/n, 28040 Madrid, Spain
Correspondence to: Ester Rojo; e-mail: firstname.lastname@example.org
Contract grant sponsor: Ministerio de Economta y Competitividad; contract grant number: CTQ2010-15742.
Published online in Wiley Online Library (wileyonlinelibrary.com)
TABLE 1. Experimental conditions and results obtained in tensile (T) and flexural (F) tests for the phenolic matrix curing. Run T ([degrees]C) t ((h) [sigma] (MPa) [epsilon] (%) T F T F 1 50 1 7.31 11.78 5.57 5.86 2 70 2.5 8.54 12.82 4.21 4.41 3 90 1 8.22 12.03 3.81 3.84 4 70 0.4 7.17 11.58 4.53 5.69 5 70 2.5 8.40 12.54 4.28 4.60 6 50 4 7.82 11.28 5.12 5.10 7 70 2.5 8.41 12.70 4.39 4.74 8 70 4.6 8.15 12.74 3.77 4.66 9 41.7 2.5 7.24 11.72 5.44 5.81 10 90 4 7.88 13.57 2.91 3.76 11 98.3 2.5 8.16 13.84 2.85 3.62 Run E (MPa) SED (J x [10.sup.-15]/ [m.sup.3]) T F T F 1 232.83 212.17 2.43 3.94 2 290.90 246.56 2.14 3.06 3 294.86 273.62 1.82 2.65 4 240.36 221.34 1.98 3.66 5 300.94 249.29 2.12 3.05 6 241.17 211.32 2.47 3.41 7 289.29 236.91 2.16 3.24 8 286.75 278.31 1.82 2.93 9 231.30 211.23 2.56 3.87 10 309.05 320.85 1.25 1.78 11 332.13 317.85 1.26 2.42 TABLE 2. Fiber treatment conditions and code names of the cellulose fiber-re in forced phenolic composites. Composites Fiber treatment Treatment conditions Cell-Ph Untreated -- Cell-1% NaOH-Ph Alkali treatments 1% NaOH, 120 min Cell--5% NaOH-Ph 5% NaOH, 120 min Cell-APS-Ph Silane treatments 2.2% silane APS, 120 min Cell-AAPS-Ph 1.5% silane AAPS, 100 min Cell-1% NaOH + APS-Ph Combined treatments 1% NaOH, 120 min + 2.2% APS, 120 min Cell-1% NaOH + AAPS-Ph 1% NaOH, 120 min + 1.5% AAPS, 100 min Cell--5% NaOH + APS-Ph 5% NaOH, 120 min + 2.2% APS, 120 min Cel 1-5% NaOH + AAPS-Ph 5% NaOH, 120 min + 1.5% AAPS, 100 min TABLE 3. Comparison between predicted and experimental values of mechanical properties for the phenolic matrix cured under optimal conditions. Tension Response Predicted Experimental [sigma] (MPa) 8.51 8.50 [epsilon] (%) 3.97 3.91 E (MPa) 297.29 282.14 SED (J/[m.sup.3]) 1.99 x [10.sup.5] 1.87 x [10.sup.5] Tension Flexion Response Dried (a) Predicted [sigma] (MPa) 13.28 [156%] 12.90 [epsilon] (%) 1.86 [48%1 4.49 E (MPa) 835.04 [296%] 265.38 SED (J/[m.sup.3]) 1.32 x [10.sup.5] [71%] 2.89 x [10.sup.5] Flexion Response Experimental Dried (a) [sigma] (MPa) 12.82 24.20 [189%] [epsilon] (%) 4.93 2.31 [47%1 E (MPa) 262.06 689.46 [263%] SED (J/[m.sup.3]) 2.79 x [10.sup.5] 2.53 x [10.sup.5] [91%] (a) Mechanical properties of the phenolic matrix after drying Values in parentheses represent the mechanical properties of dried matrices referred to those of the undried material (100%) TABLE 4. XRD position (20), d-spacing, and relative intercalation (RI) of cellulose fiber, phenolic matrix, and composites. Sample 20 ([degrees]) d-spacing (nm) RI (%) Cell 12.6810 69.8082 -- Unreinforced -- -- -- Cell-Ph 11.5597 76.5530 9.66 Cell--l% NaOH-Ph 11.5601 76.5503 9.66 Cell--5% NaOH-Ph 11.5594 76.5546 9.66 Cell-APS-Ph 11.5161 76.8421 10.08 Cell-AAPS-Ph 11.4909 77.0098 10.32 Cell--1% NaOH + APS-Ph 11.5503 76.6151 9.75
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|Author:||Rojo, Ester; Oliet, Mercedes; Alonso, M. Virginia; Del Saz-Orozco, Belen; Rodriguez, Francisco|
|Publication:||Polymer Engineering and Science|
|Date:||Oct 1, 2014|
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