**3. Results and discussion**

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

Differential scanning calorimetry was performed on PP specimens of about 15 mg by means of a Mettler DSC30 calorimeter by thermal cycling in the range 0–220°C with a heating/cooling rate of 10°C/min. Melting and crystallization temperature/peak were registered. The crystallinity X was determined referring the measured melting enthalpy ∆Hi in the first heating scan to 207 J/g and the standard

X = 100 ∆Hi/207. (2)

DMTA-thermal creep specimens of WNW (stripes 20x5 mm) and BCF (15 mm length) were subjected to dynamical mechanical analysis in tensile mode by using a DMTA Mk II (Polymer Laboratories) with a dynamic deformation of 11 μm, frequency of 5 Hz, static stress between 1 and 16 MPa, and heating rate of 3°C/min in the range −50/120°C. Storage (E') and loss (E") moduli are reported as a function

Isothermal creep at 25°C was performed on samples 200 mm length (and 50 mm width for WNW fabrics) by applying for 30 mins a constant stress of 16 MPa for BCF and 3.5 MPa for WNW, and a following recovery for 30 mins at a minimum stress of 1.5 MPa for BCF and 0.05 MPa for WNW, respectively. WNW fabrics were

All diffraction images were collected in transmission using a modified Laue

at a distance from the sample equal to 8.81 cm. The sample to detector distance

(∆L/L0) (3)

with a pixel size of 43 microns

of temperature. Thermal creep (TC) was also evaluated according to Eq. (3):

where ∆L is the specimen length variation and L0 is the initial length.

length) by using a dynamometer Instron mod. 4502 with a crosshead speed of 100 mm/min. WNW fabrics were tested both in machine and in cross direction. Single fibers of 10–20 mm were tested with a crosshead of 5–10 mm/min.

enthalpy of the fully crystalline PP according to Eq. (2):

TC = 100<sup>∗</sup>

tested both in machine and cross direction.

camera with a removable image plate 24 × 15 cm2

**2.3 Fiber diffraction measurements**

*Fiber diffraction raw images for the PP fibers with 10 wt% kaolinite (K10). (a) As-span fibers showing smoother texture. (b) Fibers at DR = 10: the central spots radially dispersed are due to diffraction from the residual bremsstrahlung radiation in the filtered only X-ray beam. Sharper spots are produced by the PP strong* 

*fiber texture, and more continuous circles are from kaolinite diffraction.*

**94**

**Figure 2.**

### **3.1 Mechanical properties**

The effect of orientation of various polypropylene products can be firstly evaluated by the different mechanical properties; in particular modulus and maximum stress (strength) are summarized in **Figure 3**.

Dumbbell specimens exhibited a tensile modulus of 950 ± 9 MPa, yield stress of 33 ± 1 MPa, and stress and deformation at break of 15 ± 2 MPa and 73%, respectively. Mechanical properties are quite different from other products, even if molecular weight is quite similar to BCF products. In injection molding, chain alignment and solidification follow a different pattern with respect to the fiber formation during spinning, and consequently dumbbell specimens exhibit heterogeneity in macromolecular orientation with a skin effect and a disordered core structure, as well described in literature [11]. Moreover, it should be considered that the shape factor, calculated as the ratio between the perimeter and section, is lower for IM (0.8 mm<sup>−</sup><sup>1</sup> ) with respect to 160 mm<sup>−</sup><sup>1</sup> for WNW, 80 mm<sup>−</sup><sup>1</sup> for BCF, and in the range of 8–31 mm<sup>−</sup><sup>1</sup> for FS depending on the drawing. The efficiency of chain orientation

#### **Figure 3.**

*Comparison of mechanical properties (tensile modulus and strength) of various oriented products, such as WNW, undrawn fibers from WNW, BCF, and drawn fibers with DR = 10; selected data of PP fibers or nanocomposite fibers with 10 (K10), 20 (K20), and 30% (K30) of kaolinite. Data of IM sample are also shown for comparison.*

**Figure 4.** *Comparison of stress-strain curves of BCF filament and WNW fabrics tested in MD and CD directions.*

is more evident in fiber-like products, where processing conditions determine a linear stretching (drawing) of polymer with elongation flow, either in spinning or in drawing.

The higher the orientation, the higher the modulus, the higher the strength, and in general the lower the strain at break. It is evident of the lower values of WNW fabrics with respect to undrawn fibers and BCF filaments. An almost linear dependence between strength and stiffness can be observed in **Figure 3**.

The highest modulus and strength have been obtained for lab-scale fiber after drawing 10 times (DR10) and relatively higher values for nanocomposites with kaolinite between 10 and 30%.

Stress-strain curves of WNW fabrics and BCF filaments are shown in **Figure 4**. The effect of fiber spinning in BCF and the direct drawing in the process with draw ratio of about 4 determines not only a stiffening of the filament but also an increase of max stress (320 MPa) and a relatively high deformation at failure (125%). More peculiar is the orientation in WNW fabric where longitudinal or machine direction and transversal or cross direction determine a different mechanical behavior. Anisotropy of WNW products is very common and usually decreases with the increase of surface density.

This effect is more evident in the creep curve shown in **Figure 5**, where CD sample exhibited not only a higher deformation after 30 mins creep with respect to MD sample (3.1 vs. 2.0%) but also a higher residual after 30 mins recovery (about 0.8 vs. 0.3%).

Correspondingly a higher storage modulus was found for MD sample with respect to CD sample, as shown in **Figure 6**. On the other hand the peak of the polymer glass transition temperature (Tg) remains localized at about 0°C.

The different thermal creep (at 3.5 MPa) of WNW with respect to BCF is reported in **Figure 7**, and it decreases in the order WNW-CD > WNW-MD > BCF. In all cases the deformation starts at about 30°C after the glass transition interval depicted by the loss modulus peak in the interval −15/30°C. As expected, the higher the orientation, the higher the storage modulus, the lower the loss modulus, and consequently the lower the thermal creep.

**Figures 8** and **9** show DMTA data of BCF filaments. Storage modulus and thermal creep directly depend on the applied stress.

Moreover, it is evident that static stress of 1 or 7 MPa determines a thermal creep starting from about 30°C, as in the case of WNW with 3.5 MPa (see **Figure 7**). On the other hand, at higher stress of 16 MPa, thermal creep starts at 0°C that correspond to the Tg, measured at the maximum of loss modulus peak.

**97**

**Figure 7.**

*mechanical analysis with static stress of 3.5 MPa.*

**Figure 5.**

**Figure 6.**

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene…*

*Comparison of creep curves of WNW fabrics tested in MD and CD directions at 25°C with an applied stress of* 

*Dynamic mechanical analysis of WNW fabrics performed at a static stress of 3.5 MPa. Storage modulus ( )*

*Thermal creep comparison of WNW fabrics (MD and CD) and BCF filaments as measured by dynamic* 

 *and loss modulus ( ) of MD and CD samples are compared.*

*DOI: http://dx.doi.org/10.5772/intechopen.85554*

*3.5 MPa for 30 mins, followed by 30 mins recovery.*

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene… DOI: http://dx.doi.org/10.5772/intechopen.85554*

#### **Figure 5.**

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

is more evident in fiber-like products, where processing conditions determine a linear stretching (drawing) of polymer with elongation flow, either in spinning or

*Comparison of stress-strain curves of BCF filament and WNW fabrics tested in MD and CD directions.*

dence between strength and stiffness can be observed in **Figure 3**.

The higher the orientation, the higher the modulus, the higher the strength, and in general the lower the strain at break. It is evident of the lower values of WNW fabrics with respect to undrawn fibers and BCF filaments. An almost linear depen-

The highest modulus and strength have been obtained for lab-scale fiber after drawing 10 times (DR10) and relatively higher values for nanocomposites with

Stress-strain curves of WNW fabrics and BCF filaments are shown in **Figure 4**. The effect of fiber spinning in BCF and the direct drawing in the process with draw ratio of about 4 determines not only a stiffening of the filament but also an increase of max stress (320 MPa) and a relatively high deformation at failure (125%). More peculiar is the orientation in WNW fabric where longitudinal or machine direction and transversal or cross direction determine a different mechanical behavior. Anisotropy of WNW products is very common and usually decreases with the increase of surface density. This effect is more evident in the creep curve shown in **Figure 5**, where CD sample exhibited not only a higher deformation after 30 mins creep with respect to MD sample (3.1 vs. 2.0%) but also a higher residual after 30 mins recovery (about

Correspondingly a higher storage modulus was found for MD sample with respect to CD sample, as shown in **Figure 6**. On the other hand the peak of the polymer glass transition temperature (Tg) remains localized at about 0°C. The different thermal creep (at 3.5 MPa) of WNW with respect to BCF is reported in **Figure 7**, and it decreases in the order WNW-CD > WNW-MD > BCF. In all cases the deformation starts at about 30°C after the glass transition interval depicted by the loss modulus peak in the interval −15/30°C. As expected, the higher the orientation, the higher the storage modulus, the lower the loss modulus, and

**Figures 8** and **9** show DMTA data of BCF filaments. Storage modulus and

Moreover, it is evident that static stress of 1 or 7 MPa determines a thermal creep starting from about 30°C, as in the case of WNW with 3.5 MPa (see **Figure 7**). On the other hand, at higher stress of 16 MPa, thermal creep starts at 0°C that corre-

**96**

in drawing.

**Figure 4.**

0.8 vs. 0.3%).

kaolinite between 10 and 30%.

consequently the lower the thermal creep.

thermal creep directly depend on the applied stress.

spond to the Tg, measured at the maximum of loss modulus peak.

*Comparison of creep curves of WNW fabrics tested in MD and CD directions at 25°C with an applied stress of 3.5 MPa for 30 mins, followed by 30 mins recovery.*

#### **Figure 6.**

*Dynamic mechanical analysis of WNW fabrics performed at a static stress of 3.5 MPa. Storage modulus ( ) and loss modulus ( ) of MD and CD samples are compared.*

#### **Figure 7.**

*Thermal creep comparison of WNW fabrics (MD and CD) and BCF filaments as measured by dynamic mechanical analysis with static stress of 3.5 MPa.*

**Figure 8.**

*Dynamic mechanical analysis of BCF filaments performed at different static stresses of 1, 7, and 16 MPa. Storage modulus ( ) and loss modulus ( ).*

**Figure 9.**

*Comparison of thermal creep measured in dynamic mechanical analysis of BCF filaments performed at different static stresses of 1, 7, and 16 MPa.*

Creep curves of BCF filaments at 16 and 78 MPa are compared in **Figure 10**. Experimental data have been interpolated by using a simple exponential model, as previously described [16]. The two parameter K and n formally represent the intensity and the rate of the creep, and they are directly dependent on the applied stress.

Thermal analysis (DSC data) could be useful for a preliminary evaluation of samples. The first heating scan is representative of the products, whereas the cooling step and the second heating scan give information on the polymer. Crystallinity (endothermal peak), the quality of the peak from melting peak and crystallizability of polymer during the cooling step, usually depends on molecular weight. For IM sample melting temperature of 169.2°C and melting enthalpy of 83.2 J/g (crystallinity of 40.2%) were determined. **Tables 2** and **3** show selected data of WNW and BCF samples, respectively. DSC data of various WNW fabrics before and after testing are almost similar.

Thermal analysis of BCF original filaments and after creep or mechanical experiments revealed an interesting result of increased crystallization, in dependence on the well-known phenomenon described as "stress-induced crystallization" [17] or "orientation-induced crystallization" [18] that occurs either in fiber processing or

**99**

**Table 3.**

**Figure 10.**

**Table 2.**

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene…*

*Comparison of creep curves of BCF filament tested at 25.0°C with different high applied stresses of 11 or* 

**ΔH1 [J/g]**

**WNW First Scan Cooling Second Scan**

As received 175.5 84.6 40.5 121.1 95.2 164.2 96.2 MD after creep 3.5 MPa 171.6 84.1 40.2 121.5 96.3 163.4 96.6 CD after creep 3.5 MPa 173.7 84.3 40.3 121.8 95.1 163.6 96.7 MD after failure 170.1 86.2 41.2 121.8 95.7 163.5 95.8 CD after failure 173.1 82.5 39.9 120.1 94.1 165.1 95.2

*Thermal results of WNW fabrics (melting and crystallization temperature; melting and crystallization* 

**ΔH1 [J/g]**

**BCF First Scan Cooling Second Scan**

As received 175.5 74.6 35.7 114.4 90.9 166.6 91.3 After creep 16 MPa 172.9 74.0 35.4 115.1 89.4 165.5 90.7 After creep 78 MPa 172.9 79.4 38.0 115.4 89.5 165.5 91.5 After failure 174.0 108.3 51.8 115.4 89.1 166.0 90.5 *Melting and crystallization temperature; melting and crystallization enthalpy in the three DSC scans. Crystallinity is* 

**X [%]**

**Tc [°C]**

**ΔHc [J/g]**

**X [%]**

**Tc [°C]**

**ΔHc [J/g]**

**Tm2 [°C]**

**Tm2 [°C]**

**ΔH2 [J/g]**

**ΔH2 [J/g]**

in fiber testing. The residual deformation after creep and recovery was found at 0.7 and 7.5% after loading 16 or 78 MPa, respectively, and correspondently crystallinity of 35 and 38% was measured. It should be noted that crystallinity of BCF filaments

increased up to 52% after the failure (strain failure at 125%).

*DOI: http://dx.doi.org/10.5772/intechopen.85554*

*78 MPa for 30 mins, followed by 30 mins of recovery.*

*Crystallinity is calculated according to Eq. (2).*

*enthalpy in the three DSC scans).*

*calculated according to Eq. (2).*

*Thermal results of BCF filaments.*

**Tm1 [°C]**

**Tm1 [°C]** *Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene… DOI: http://dx.doi.org/10.5772/intechopen.85554*

#### **Figure 10.**

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

Creep curves of BCF filaments at 16 and 78 MPa are compared in **Figure 10**. Experimental data have been interpolated by using a simple exponential model, as previously described [16]. The two parameter K and n formally represent the intensity and the rate of the creep, and they are directly dependent on the

*Comparison of thermal creep measured in dynamic mechanical analysis of BCF filaments performed at* 

*Dynamic mechanical analysis of BCF filaments performed at different static stresses of 1, 7, and 16 MPa.* 

Thermal analysis (DSC data) could be useful for a preliminary evaluation of samples. The first heating scan is representative of the products, whereas the cooling step and the second heating scan give information on the polymer. Crystallinity (endothermal peak), the quality of the peak from melting peak and crystallizability of polymer during the cooling step, usually depends on molecular weight. For IM sample melting temperature of 169.2°C and melting enthalpy of 83.2 J/g (crystallinity of 40.2%) were determined. **Tables 2** and **3** show selected data of WNW and BCF samples, respectively. DSC data of various WNW fabrics before and after

Thermal analysis of BCF original filaments and after creep or mechanical experiments revealed an interesting result of increased crystallization, in dependence on the well-known phenomenon described as "stress-induced crystallization" [17] or "orientation-induced crystallization" [18] that occurs either in fiber processing or

**98**

applied stress.

**Figure 9.**

**Figure 8.**

testing are almost similar.

*different static stresses of 1, 7, and 16 MPa.*

*Storage modulus ( ) and loss modulus ( ).*

*Comparison of creep curves of BCF filament tested at 25.0°C with different high applied stresses of 11 or 78 MPa for 30 mins, followed by 30 mins of recovery.*


#### **Table 2.**

*Thermal results of WNW fabrics (melting and crystallization temperature; melting and crystallization enthalpy in the three DSC scans).*


*Melting and crystallization temperature; melting and crystallization enthalpy in the three DSC scans. Crystallinity is calculated according to Eq. (2).*

#### **Table 3.**

*Thermal results of BCF filaments.*

in fiber testing. The residual deformation after creep and recovery was found at 0.7 and 7.5% after loading 16 or 78 MPa, respectively, and correspondently crystallinity of 35 and 38% was measured. It should be noted that crystallinity of BCF filaments increased up to 52% after the failure (strain failure at 125%).

#### **Figure 11.**

*Comparison of max stress and strain at failure of PP fibers in the two steps of processing, i.e., spinning and drawing. Interpolation curve following the equation proposed in literature [17].*

#### **3.2 Case of highly oriented products (fibers)**

Single fibers at different levels of drawing were prepared and characterized. **Figure 11** shows the relationship between the max stress and deformation at break that could be interpolated according to the criteria of independence of a total maximum orientation or the network deformation concept [17]. It could be assumed as a direct combination of the orientation obtained in fiber spinning, the extension of the fiber during fiber drawing, and the strain deformation during tensile testing [17]. Experimental results of PP fibers have interpolated separately, distinguishing the spinning step and the drawing step.

The calculated values of 827 and 1225 MPa formally represent the maximum attainable strength in fiber spinning and fiber drawing. It is well evident that 1225 MPa is underestimated, suggesting that various and different phenomena occur during fiber production with different roles and consequences on the ultimate properties.

Moreover, the effect of the filler was also observed and evaluated in drawing, by comparing the modulus and the strength of the fiber as function of draw ratio, as shown in **Figure 12**. The higher the draw ratio, the higher the orientation, and the higher the modulus and the strength, for both PP fiber and composite fibers.

The maximum modulus and strength have been obtained after drawing at draw ratio of DR 15 for both iPP and composite fibers, with values in the range of 8–9 GPa and 900–990 MPa, respectively.

#### **3.3 Texture analysis**

Fiber diffraction using transmission images is a powerful technique to analyze polymers especially in fiber or textile form. From a just one diffraction image, it may be possible to get crystal structure information [19, 20], texture [21, 22], and also microstructural features [21]. A strong texture and crystallization may help the crystal structure solution and refinement the same way as the texture is used for crystal structure solutions [22]. The major problem in the quantitative analysis of these transmission images is to account for the texture in a correct way. Simple fiber

**101**

**Figure 12.**

that was our principal goal.

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene…*

textures can be modeled easily as proven by Ran et al. [23], and in-line analyses can be carried out to monitor the crystallizable behavior of polypropylene fibers. But only one attempt has been made to obtain some rough quantitative information on the texture by Jin et al. [24] for the polypropylene fibers. In this paper we will show a procedure to perform a global diffraction analysis from which we can obtain simultaneously all information from the crystal structure to the orientation distribution function (ODF). The ODF is a function describing the volumetric amount of material with a specific crystallographic orientation in one direction. The procedure is based on the Rietveld texture analysis [25, 26] but using transmission images [20, 21]. To briefly recall the general methodology, the 2D images collected in transmission are transformed in spectra sampling the image in radial slices each one covering a certain diffraction cone angle [27–29]. All the spectra are refined at once in the Rietveld refinement program MAUD [30] using in addition to the crystal structure and size-strain model a texture model to obtain the ODF. In the case of samples in fiber form, the more suitable model is the standard functions [20, 31] as it provides a sufficiently flexible way to represent the ODF with few refined parameters. Instead for smoother texture, like in the case of WNW samples, which we could not model with an ideal standard function, the more flexible spherical harmonic method [25, 31–33] has been used but in the exponential form to ensure a positive function.

*Modulus and strength of PP fibers and nanocomposites with 10 and 20% of kaolinite as function of draw ratio.*

Only isotactic polypropylene (iPP) was found in our samples, and for the fiber diffraction images fitting and refinement, we started from the crystal structure determined by Natta et al. [34, 35]. In order to reduce the number of degrees of freedom in our model, we used a bond and angle restraint function for the iPP. This was sufficient to drive our analysis to a unique solution and safely obtain the ODF

**Figure 13** reports the fitting for the BCF as-span sample. The same analysis procedure was applied also to the strained fibers after creep/recovery at 78 MPa with residual deformation of 7.5% (as shown in **Figure 10**) and after thermal creep in DMTA at 16 MPa with residual deformation of 30% (see **Figure 9**). The experimental diffraction image on the left was transformed by unrolling around the

*DOI: http://dx.doi.org/10.5772/intechopen.85554*

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene… DOI: http://dx.doi.org/10.5772/intechopen.85554*

**Figure 12.** *Modulus and strength of PP fibers and nanocomposites with 10 and 20% of kaolinite as function of draw ratio.*

textures can be modeled easily as proven by Ran et al. [23], and in-line analyses can be carried out to monitor the crystallizable behavior of polypropylene fibers. But only one attempt has been made to obtain some rough quantitative information on the texture by Jin et al. [24] for the polypropylene fibers. In this paper we will show a procedure to perform a global diffraction analysis from which we can obtain simultaneously all information from the crystal structure to the orientation distribution function (ODF). The ODF is a function describing the volumetric amount of material with a specific crystallographic orientation in one direction. The procedure is based on the Rietveld texture analysis [25, 26] but using transmission images [20, 21].

To briefly recall the general methodology, the 2D images collected in transmission are transformed in spectra sampling the image in radial slices each one covering a certain diffraction cone angle [27–29]. All the spectra are refined at once in the Rietveld refinement program MAUD [30] using in addition to the crystal structure and size-strain model a texture model to obtain the ODF. In the case of samples in fiber form, the more suitable model is the standard functions [20, 31] as it provides a sufficiently flexible way to represent the ODF with few refined parameters. Instead for smoother texture, like in the case of WNW samples, which we could not model with an ideal standard function, the more flexible spherical harmonic method [25, 31–33] has been used but in the exponential form to ensure a positive function.

Only isotactic polypropylene (iPP) was found in our samples, and for the fiber diffraction images fitting and refinement, we started from the crystal structure determined by Natta et al. [34, 35]. In order to reduce the number of degrees of freedom in our model, we used a bond and angle restraint function for the iPP. This was sufficient to drive our analysis to a unique solution and safely obtain the ODF that was our principal goal.

**Figure 13** reports the fitting for the BCF as-span sample. The same analysis procedure was applied also to the strained fibers after creep/recovery at 78 MPa with residual deformation of 7.5% (as shown in **Figure 10**) and after thermal creep in DMTA at 16 MPa with residual deformation of 30% (see **Figure 9**). The experimental diffraction image on the left was transformed by unrolling around the

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

Single fibers at different levels of drawing were prepared and characterized. **Figure 11** shows the relationship between the max stress and deformation at break that could be interpolated according to the criteria of independence of a total maximum orientation or the network deformation concept [17]. It could be assumed as a direct combination of the orientation obtained in fiber spinning, the extension of the fiber during fiber drawing, and the strain deformation during tensile testing [17]. Experimental results of PP fibers have interpolated separately, distinguishing

*Comparison of max stress and strain at failure of PP fibers in the two steps of processing, i.e., spinning and* 

The calculated values of 827 and 1225 MPa formally represent the maximum attainable strength in fiber spinning and fiber drawing. It is well evident that 1225 MPa is underestimated, suggesting that various and different phenomena occur during fiber

Moreover, the effect of the filler was also observed and evaluated in drawing, by comparing the modulus and the strength of the fiber as function of draw ratio, as shown in **Figure 12**. The higher the draw ratio, the higher the orientation, and the higher the modulus and the strength, for both PP fiber and composite fibers. The maximum modulus and strength have been obtained after drawing at draw ratio of DR 15 for both iPP and composite fibers, with values in the range of

Fiber diffraction using transmission images is a powerful technique to analyze polymers especially in fiber or textile form. From a just one diffraction image, it may be possible to get crystal structure information [19, 20], texture [21, 22], and also microstructural features [21]. A strong texture and crystallization may help the crystal structure solution and refinement the same way as the texture is used for crystal structure solutions [22]. The major problem in the quantitative analysis of these transmission images is to account for the texture in a correct way. Simple fiber

production with different roles and consequences on the ultimate properties.

**3.2 Case of highly oriented products (fibers)**

*drawing. Interpolation curve following the equation proposed in literature [17].*

the spinning step and the drawing step.

8–9 GPa and 900–990 MPa, respectively.

**3.3 Texture analysis**

**Figure 11.**

**100**

**Figure 13.**

*Original fiber diffraction image for the BCF as-span on the left. Unrolled and fitted data at the end of the analysis on the right. The lower part of the right image (spectrum number from 0 to 143) contains the unrolled experimental data, on the upper part (spectrum number 144 to 287), the calculated patterns. The matching of the two parts indicates a good fitting and correct model.*

diffraction center (the white hole) in steps of 2.5° in order to get 144 radial diffraction spectra that were fitted by the program MAUD. **Figure 13**, on the right, shows the result of the fitting. The calculated patterns in the top part well reproduce the unrolled experimental spectra in the bottom. The longer spots at low angles, and visible in the image on the right as radial short strings, are due to the diffraction of the bremsstrahlung that was also included in the pattern modeling. In the unrolled map on the right, the presence of sharper spots vertically means sharper texture. Instead, sharper spots horizontally correspond to a better crystallization of the fibers/compound.

In **Figure 14** we have recalculated some of the pole figures from the ODFs obtained for the three BCF samples. Pole figures, being 2D, are somehow more convenient to visualize the texture characteristics. All three different fibers show the same kind of texture: the fiber axis is parallel to the normal to the pole figures, and a perfect fiber symmetry was found for the (100) axis. For the texture analysis, we were obliged to use the monoclinic c-setting for the iPP instead of the more common b-setting. The usual (001) fiber axis becomes the (100) in our case, because of the different cell conventions used. The texture sharpness does not change significantly between the as-span and strained fibers at room temperature, but there is a strong increase in fiber alignment after thermal creep. In fact the fiber spread obtained from the fitting was 15 ± 1° for the as-span and 17 ± 1° for the room temperature creep but becomes 9 ± 1° for the fibers after creep in DMTA up to 120°C (see **Figure 9**). Also the mean crystallite sizes, as measured from the analysis, increase from 14 to 19 nm with the thermal creep, but it is not affected by the creep at room temperature. These findings are in agreement with the residual deformation of BCF filament, i.e., 30% after thermal creep (**Figure 9**) and 7.5% after creep/ recovery (**Figure 10**).

For the WNM samples, for which we show only one experimental diffraction image and its fitting in **Figure 15**, the results of the texture analysis for the three samples, as received and after creep in MD and CD directions, are shown in **Figure 16**. The texture is weaker than in the previous BCF case and it is not a fiber one. The machine and cross directions show a slightly different fiber alignment that gives result to a different texture when subjected to creep in their respective direction. The texture sharpness increases a bit and more for the machine direction, as it was showing a more favorable alignment of the fibers from the beginning.

**103**

**Figure 15.**

**Figure 14.**

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene…*

Finally, we analyzed in fiber diffraction also one sample containing kaolinite. The analysis was more difficult in this case as two textured phases are present. Again, kaolinite shows a high density of modulated planar defects and strong

*Original fiber diffraction image (left) and unrolled and fitted patterns (right) for the WNW sample. The texture is much smoother with respect to the drawn fiber samples. With respect to* **Figure 13***, the uncomplete 2Θ*

*Recalculated pole figures for BCF fibers (from top): as-span, after creep at 78 MPa (room temperature) and after thermal creep at 16 MPa. Only after creep at high temperature we notice an increase in the crystallographic* 

*alignment of the fibers. The fiber sample direction corresponds to the normal to the pole figures.*

compression in the plane normal to the fibers axis.

*range has been cut to enhance low-angle features.*

*DOI: http://dx.doi.org/10.5772/intechopen.85554*

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene… DOI: http://dx.doi.org/10.5772/intechopen.85554*

#### **Figure 14.**

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

diffraction center (the white hole) in steps of 2.5° in order to get 144 radial diffraction spectra that were fitted by the program MAUD. **Figure 13**, on the right, shows the result of the fitting. The calculated patterns in the top part well reproduce the unrolled experimental spectra in the bottom. The longer spots at low angles, and visible in the image on the right as radial short strings, are due to the diffraction of the bremsstrahlung that was also included in the pattern modeling. In the unrolled map on the right, the presence of sharper spots vertically means sharper texture. Instead, sharper spots horizontally correspond to a better crystallization of the

*Original fiber diffraction image for the BCF as-span on the left. Unrolled and fitted data at the end of the analysis on the right. The lower part of the right image (spectrum number from 0 to 143) contains the unrolled experimental data, on the upper part (spectrum number 144 to 287), the calculated patterns. The matching of* 

In **Figure 14** we have recalculated some of the pole figures from the ODFs obtained for the three BCF samples. Pole figures, being 2D, are somehow more convenient to visualize the texture characteristics. All three different fibers show the same kind of texture: the fiber axis is parallel to the normal to the pole figures, and a perfect fiber symmetry was found for the (100) axis. For the texture analysis, we were obliged to use the monoclinic c-setting for the iPP instead of the more common b-setting. The usual (001) fiber axis becomes the (100) in our case, because of the different cell conventions used. The texture sharpness does not change significantly between the as-span and strained fibers at room temperature, but there is a strong increase in fiber alignment after thermal creep. In fact the fiber spread obtained from the fitting was 15 ± 1° for the as-span and 17 ± 1° for the room temperature creep but becomes 9 ± 1° for the fibers after creep in DMTA up to 120°C (see **Figure 9**). Also the mean crystallite sizes, as measured from the analysis, increase from 14 to 19 nm with the thermal creep, but it is not affected by the creep at room temperature. These findings are in agreement with the residual deformation of BCF filament, i.e., 30% after thermal creep (**Figure 9**) and 7.5% after creep/

For the WNM samples, for which we show only one experimental diffraction image and its fitting in **Figure 15**, the results of the texture analysis for the three samples, as received and after creep in MD and CD directions, are shown in **Figure 16**. The texture is weaker than in the previous BCF case and it is not a fiber one. The machine and cross directions show a slightly different fiber alignment that gives result to a different texture when subjected to creep in their respective direction. The texture sharpness increases a bit and more for the machine direction, as it was showing a more favorable

**102**

fibers/compound.

*the two parts indicates a good fitting and correct model.*

**Figure 13.**

recovery (**Figure 10**).

alignment of the fibers from the beginning.

*Recalculated pole figures for BCF fibers (from top): as-span, after creep at 78 MPa (room temperature) and after thermal creep at 16 MPa. Only after creep at high temperature we notice an increase in the crystallographic alignment of the fibers. The fiber sample direction corresponds to the normal to the pole figures.*

#### **Figure 15.**

*Original fiber diffraction image (left) and unrolled and fitted patterns (right) for the WNW sample. The texture is much smoother with respect to the drawn fiber samples. With respect to* **Figure 13***, the uncomplete 2Θ range has been cut to enhance low-angle features.*

Finally, we analyzed in fiber diffraction also one sample containing kaolinite. The analysis was more difficult in this case as two textured phases are present. Again, kaolinite shows a high density of modulated planar defects and strong compression in the plane normal to the fibers axis.

**Figure 16.**

*Recalculated pole figures for WNW samples (from top): as prepared, after creep in machine direction, and after creep in cross direction. The pole figures' horizontal direction corresponds to the normal to the WNW in-plane tissue. The creep direction is normal to the pole figures for MD and vertical for CD.*

#### **Figure 17.**

*Rietveld fitting of the unrolled diffraction image for the sample K10. The spectrum number corresponds to a pattern integrated radially every 2° in the circumferential direction starting from the horizontal plane. There are 180 experimental integrated patterns in the lower part and 180 recalculated patterns in the upper part. The waving vertical lines correspond to the kaolinite diffraction interplanar spacings in compression in the plane perpendicular to the fibers.*

**105**

inserted (10 wt%).

**Figure 18.**

*fiber direction is normal to the pole figures.*

**4. Conclusions**

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene…*

To correctly model the unrolled diffraction patterns, as reported in **Figure 17**, we had to derive a modulated planar defect model for kaolinite starting from a previous turbostratic model [36]. This modulated defect structure has been tested by analyzing

*Recalculated pole figures of iPP (top) and kaolinite (bottom) resulting from the Rietveld texture fitting. The* 

The fitting of the fiber diffraction, for the iPP with 10 wt% kaolinite at DR = 10,

From the crystallization point of view, the iPP is similar to the BCF after thermal creep (17 nm for the mean crystallite sizes), and kaolinite is arranged in packets of

**Figure 18** is showing the recalculated pole figures for the iPP and kaolinite. The fiber spread of the iPP corresponds to 7.3 ± 0.5°, and so kaolinite contributes to the alignment of the fibers reaching a texture even sharper than the BCF after thermal creep. From the kaolinite pole figures, we can deduce that the (001) basal plane is

Finally, from the analysis of the fiber diffraction image, we can measure also the amount of kaolinite, and thanks to the proper modeling, including texture and strain effects, the refined amount was 10.7 wt% that is very close to the amount

It can be concluded that polymer processing directly affects mechanical properties of iPP fiber-like products, as observed in the case of WNW fabrics and fibers, both BCF and monofilaments. The higher the orientation, the higher the modulus and the strength, the lower the deformation at break, and the higher the creep resistance. No particular variation of crystallinity of specimen before and after creep and fracture test was detected by DSC analysis, except for BCF after failure.

the powder diffraction pattern of the same kaolinite used as filler in the fibers.

is reported in **Figure 17**. The iPP texture and microstructure characteristics are similar to the BCF, especially the one after thermal creep, and by difference with **Figure 13**, the reader may recognize which ones are the diffraction spots and lines of kaolinite. Kaolinite shows a high compression state perpendicular to the fiber axis corresponding to about 0.8% equivalent elastic strain. This extremely high deformation is probably induced by the fiber drawing and the intercalation with kaolinite. By the strong texture, the deformation mainly belongs to the (001) kaolinite basal plane

(perpendicular to the plane) that is also the intercalation and faulted plane.

distributed normal to the fiber axis but in a perfect fiber texture.

about 50 nm perpendicular to the basal plane.

*DOI: http://dx.doi.org/10.5772/intechopen.85554*

*Effect of Processing and Orientation on Structural and Mechanical Properties of Polypropylene… DOI: http://dx.doi.org/10.5772/intechopen.85554*

**Figure 18.**

*Polypropylene - Polymerization and Characterization of Mechanical and Thermal Properties*

*Rietveld fitting of the unrolled diffraction image for the sample K10. The spectrum number corresponds to a pattern integrated radially every 2° in the circumferential direction starting from the horizontal plane. There are 180 experimental integrated patterns in the lower part and 180 recalculated patterns in the upper part. The waving vertical lines correspond to the kaolinite diffraction interplanar spacings in compression in the plane* 

*Recalculated pole figures for WNW samples (from top): as prepared, after creep in machine direction, and after creep in cross direction. The pole figures' horizontal direction corresponds to the normal to the WNW* 

*in-plane tissue. The creep direction is normal to the pole figures for MD and vertical for CD.*

**104**

**Figure 17.**

**Figure 16.**

*perpendicular to the fibers.*

*Recalculated pole figures of iPP (top) and kaolinite (bottom) resulting from the Rietveld texture fitting. The fiber direction is normal to the pole figures.*

To correctly model the unrolled diffraction patterns, as reported in **Figure 17**, we had to derive a modulated planar defect model for kaolinite starting from a previous turbostratic model [36]. This modulated defect structure has been tested by analyzing the powder diffraction pattern of the same kaolinite used as filler in the fibers.

The fitting of the fiber diffraction, for the iPP with 10 wt% kaolinite at DR = 10, is reported in **Figure 17**. The iPP texture and microstructure characteristics are similar to the BCF, especially the one after thermal creep, and by difference with **Figure 13**, the reader may recognize which ones are the diffraction spots and lines of kaolinite. Kaolinite shows a high compression state perpendicular to the fiber axis corresponding to about 0.8% equivalent elastic strain. This extremely high deformation is probably induced by the fiber drawing and the intercalation with kaolinite. By the strong texture, the deformation mainly belongs to the (001) kaolinite basal plane (perpendicular to the plane) that is also the intercalation and faulted plane.

From the crystallization point of view, the iPP is similar to the BCF after thermal creep (17 nm for the mean crystallite sizes), and kaolinite is arranged in packets of about 50 nm perpendicular to the basal plane.

**Figure 18** is showing the recalculated pole figures for the iPP and kaolinite. The fiber spread of the iPP corresponds to 7.3 ± 0.5°, and so kaolinite contributes to the alignment of the fibers reaching a texture even sharper than the BCF after thermal creep. From the kaolinite pole figures, we can deduce that the (001) basal plane is distributed normal to the fiber axis but in a perfect fiber texture.

Finally, from the analysis of the fiber diffraction image, we can measure also the amount of kaolinite, and thanks to the proper modeling, including texture and strain effects, the refined amount was 10.7 wt% that is very close to the amount inserted (10 wt%).

## **4. Conclusions**

It can be concluded that polymer processing directly affects mechanical properties of iPP fiber-like products, as observed in the case of WNW fabrics and fibers, both BCF and monofilaments. The higher the orientation, the higher the modulus and the strength, the lower the deformation at break, and the higher the creep resistance. No particular variation of crystallinity of specimen before and after creep and fracture test was detected by DSC analysis, except for BCF after failure.

In the case of composite fibers, the improvement in mechanical properties of monofilament is mainly dependent on the fiber drawing, whereas only a marginal contribute of kaolinite content has been observed. For instance, the processing-drawing with draw ratio of 10 produced monofilaments of iPP and kaolinite composite with modulus and strength in the range of 5.5–6.5 GPa and of 770–870 MPa, respectively.

From the texture analyzed by X-ray, we can notice that the fiber alignment is only affected by the creep at high temperature. Stretching the fibers at room texture has a negligible effect on the texture. The situation may be different for the WNW, where the fibers have a higher mobility and the stress can change their orientation, also at lower temperatures.

In addition, from this work we can conclude that the mechanical properties are highly correlated to the texture and reverse. The fibers after thermal creep, which are showing the higher residual deformation, show also a higher increase in the texture.
