*v***i (m/s) Perforation State in the Front and Back** 433.3 433.3 433.3 *Polymers* **2022**, *14*, x FOR PEER REVIEW 10 of 22 433.3

Figure 7 shows the state of a stripped brass jacket and lead filler. Figure 7a,c shows the front and back views of the perforation on the front sheet at the velocity of 524.8 m/s, and Figure 7b,d shows the front and back views of the perforation at the velocity of 501.1 m/s. The ablative phenomenon can be observed, especially on the back view of the front sheet, resulting from a severe interaction during the penetration. Figure 7 shows the state of a stripped brass jacket and lead filler. Figure 7a,c shows the front and back views of the perforation on the front sheet at the velocity of 524.8 m/s, and Figure 7b,d shows the front and back views of the perforation at the velocity of 501.1 m/s. The ablative phenomenon can be observed, especially on the back view of the front sheet, resulting from a severe interaction during the penetration. Figure 7 shows the state of a stripped brass jacket and lead filler. Figure 7a,c shows the front and back views of the perforation on the front sheet at the velocity of 524.8 m/s, and Figure 7b,d shows the front and back views of the perforation at the velocity of 501.1 m/s. The ablative phenomenon can be observed, especially on the back view of the front sheet, resulting from a severe interaction during the penetration. Figure 7 shows the state of a stripped brass jacket and lead filler. Figure 7a,c shows the front and back views of the perforation on the front sheet at the velocity of 524.8 m/s, and Figure 7b,d shows the front and back views of the perforation at the velocity of 501.1 m/s. The ablative phenomenon can be observed, especially on the back view of the front sheet, resulting from a severe interaction during the penetration.

(**c**) 524.8 m/s (**d**) 501.1 m/s

**Figure 7.** The striped brass jacket and lead filler in the perforation of the front sheet at different

**Figure 7.** The striped brass jacket and lead filler in the perforation of the front sheet at different

#### **Table 2.** *Cont. Polymers* **2022**, *14*, x FOR PEER REVIEW 10 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 10 of 22

Figure 8 shows the failure states of the PE laminates. As shown in Figure 8a, in the side view of the PE laminates, the obvious phenomenon of delamination can be observed. As presented in Figure 8b, a penetration cavity was formed by the ogive–nose penetrator with relatively neat cutting edges accompanied by a charring layer. In addition, the PE laminate exhibited an extent of fibrillation, and the bare bunches of fibers can be observed around the penetration hole. Figure 8 shows the failure states of the PE laminates. As shown in Figure 8a, in the side view of the PE laminates, the obvious phenomenon of delamination can be observed. As presented in Figure 8b, a penetration cavity was formed by the ogive–nose penetrator with relatively neat cutting edges accompanied by a charring layer. In addition, the PE laminate exhibited an extent of fibrillation, and the bare bunches of fibers can be observed around the penetration hole. with relatively neat cutting edges accompanied by a charring layer. In addition, the PE laminate exhibited an extent of fibrillation, and the bare bunches of fibers can be observed around the penetration hole.

Figure 8 shows the failure states of the PE laminates. As shown in Figure 8a, in the side view of the PE laminates, the obvious phenomenon of delamination can be observed. As presented in Figure 8b, a penetration cavity was formed by the ogive–nose penetrator

*Polymers* **2022**, *14*, x FOR PEER REVIEW 11 of 22

**Figure 8.** Failure states of the PE laminates. (**a**) Side view of PE laminates. (**b**) Outlet of PE laminate. **Figure 8.** Failure states of the PE laminates. (**a**) Side view of PE laminates. (**b**) Outlet of PE laminate. Figure 9 presents additional detail on the perforation results at the impact velocity of

Figure 9 presents additional detail on the perforation results at the impact velocity of 433.3 m/s, with the penetrator embedded in the armor. An obvious indentation was formed on the back of the back sheet, as shown in Figure 9a. After the back sheet was removed, the head of the steel core of the A–P projectile can be seen in Figure 9b, accompanied by an extent of fibrillation. Figure 9 presents additional detail on the perforation results at the impact velocity of 433.3 m/s, with the penetrator embedded in the armor. An obvious indentation was formed on the back of the back sheet, as shown in Figure 9a. After the back sheet was removed, the head of the steel core of the A–P projectile can be seen in Figure 9b, accompanied by an extent of fibrillation. 433.3 m/s, with the penetrator embedded in the armor. An obvious indentation was formed on the back of the back sheet, as shown in Figure 9a. After the back sheet was removed, the head of the steel core of the A–P projectile can be seen in Figure 9b, accompanied by an extent of fibrillation.

(**a**) (**b**)

**Figure 9.** Additional details for the perforation results at the impact velocity of 433.3 m/s. (**a**) Inden-

tation of back plate by penetration. (**b**) Embedded penetrator. Regarding the results of Table 2 and Figures 7–9, it can be summarized from the per-**Figure 9.** Additional details for the perforation results at the impact velocity of 433.3 m/s. (**a**) Indentation of back plate by penetration. (**b**) Embedded penetrator. **Figure 9.** Additional details for the perforation results at the impact velocity of 433.3 m/s. (**a**) Indentation of back plate by penetration. (**b**) Embedded penetrator.

foration results that: (1) The average velocities of 501.1 m/s and 433.3 m/s can be taken as the ballistic limits of the UHMWPE composite armor under the impact of the ogive–nose penetrator at 467.2 m/s. (2) Petaling, as the main dominant failure mode for both face sheets, can be observed within the range of impact velocity of 501 m/s to 1026 m/s. The Regarding the results of Table 2 and Figures 7–9, it can be summarized from the perforation results that: (1) The average velocities of 501.1 m/s and 433.3 m/s can be taken as the ballistic limits of the UHMWPE composite armor under the impact of the ogive–nose penetrator at 467.2 m/s. (2) Petaling, as the main dominant failure mode for both face Regarding the results of Table 2 and Figures 7–9, it can be summarized from the perforation results that: (1) The average velocities of 501.1 m/s and 433.3 m/s can be taken as the ballistic limits of the UHMWPE composite armor under the impact of the ogive–nose penetrator at 467.2 m/s. (2) Petaling, as the main dominant failure mode for both face

sheets, can be observed within the range of impact velocity of 501 m/s to 1026 m/s. The

sheets, can be observed within the range of impact velocity of 501 m/s to 1026 m/s. The surface of both face sheets stays relatively flat, with small overall deformation except for the protruding petal–shaped holes. Small pieces of petals accompanied by gaped rifts formed the perforation. (3) Delamination and shear failure dominate the penetration process of UHMWPE laminates. Due to the low interlaminar stiffness and strength in the PE laminate, delamination is prevalent through the panel's thickness, as can be seen in Figure 10a. (4) The charring layer on the front steel plate can be observed, and more severe ablation could be noticed at the impact velocity of around 1000 m/s. surface of both face sheets stays relatively flat, with small overall deformation except for the protruding petal–shaped holes. Small pieces of petals accompanied by gaped rifts formed the perforation. (3) Delamination and shear failure dominate the penetration process of UHMWPE laminates. Due to the low interlaminar stiffness and strength in the PE laminate, delamination is prevalent through the panel's thickness, as can be seen in Figure 10a. (4) The charring layer on the front steel plate can be observed, and more severe ablation could be noticed at the impact velocity of around 1000 m/s.

**Figure 10.** Numerical model of UHMWPE composite armor and the sharp head penetrator. (**a**) Grid model. (**b**) Side view of numerical model. (**c**) Isometric side view of numerical model. **Figure 10.** Numerical model of UHMWPE composite armor and the sharp head penetrator. (**a**) Grid model. (**b**) Side view of numerical model. (**c**) Isometric side view of numerical model.

#### **4. Numerical Simulation and Analysis**

#### *4.1. Setup of Numerical Model*

To predict the dynamic response and obtain the ballistic limit of UHMWPE composite armor under the ballistic impact of the A–P core, three–dimensional numerical models are carried out using the AUTODYN nonlinear software. The version of AUTODYN is v11.0 in the software of ANSYS 11.0, located in Nanjing, China.

As shown in Figure 10, the 3D Lagrange algorithm is adopted for all of the components in numerical simulation. The half 3D model is carried out with a mesh size of about 1.2 mm per grid. A hexahedral structured grid is used to model both the projectile and the composite armor. The numerical simulation model is composed of about 810 thousand nodes and 800 elements. On the edge of the target, fixed boundaries are used to constrain the movement of the armor. The boundary conditions are applied on the edges of both the face and back sheets. Different initial velocities are applied to the ogive–nose head penetrator to simulate the dynamic penetration behavior with different impact velocities. The material models and the parameters will be described below.

As presented in Table 3, the material models for the penetrator, face sheet, and UHMWPE laminate are listed. For steel, the shock equation of state, also called Grüneisen, is employed in conjunction with the Johnson–Cook constitutive model to simulate the dynamic response under ballistic impact. The Grüneisen EOS [27] can be used to describe how the materials interact with the shock wave and are based on Hugoniot's relation between the vs. and the *v*p, as *v*<sup>s</sup> = *c*<sup>0</sup> + *sv*p, where vs. is the shock wave velocity, *v*<sup>p</sup> is the material particle velocity, *c*<sup>0</sup> is the wave speed, and *s* is a material–related coefficient. The expression of the equation of state of Grüneisen for the compressed state is:

$$p = \frac{\rho\_0 \mathcal{C}^2 \mu \left[1 + \left(1 - \frac{\gamma\_0}{2}\right) \mu - \frac{a}{2} \mu^2\right]}{\left[1 - (\mathcal{S}\_1 - 1)\mu - \mathcal{S}\_2 \frac{\mu^2}{\mu + 1} - \mathcal{S}\_3 \frac{\mu^3}{\left(\mu + 1\right)^2}\right]} + (\gamma\_0 + a\mu)E. \tag{1}$$

**Table 3.** Material models used in numerical simulation.


In the expanded state,

$$p = \rho\_0 \mathbb{C}^2 \mu + (\gamma\_0 + a\mu)E \tag{2}$$

where *C* is the intercept of the velocity curve between the shock wave and particle; *S*1, *S*2, and *S*<sup>3</sup> represent the slope of the *v*<sup>s</sup> − *v*<sup>p</sup> curve; *γ*<sup>0</sup> is the coefficient of the Grüneisen; *a* is the one–order correction of *γ*0. *µ* = *ρ*/*ρ*<sup>0</sup> − 1 is a non–dimensional coefficient based on initial and instantaneous material densities. The parameters of the Grüneisen equation of state are listed in Table 4.

**Table 4.** EOS parameters of S-7 and Q235.


The Johnson–Cook model [28,29] incorporates the effect of strain rate–dependent work hardening and thermal softening, which is given by:

$$\sigma = (A + B\varepsilon^n) \left( 1 + \mathcal{C} \ln \frac{\dot{\varepsilon}}{\dot{\varepsilon}\_0} \right) (1 - T^{\*m}) \tag{3}$$

where *ε* is the plastic strain, and the temperature factor is expressed as:

$$T^\* = \frac{T - T\_r}{T\_m - T\_r} \tag{4}$$

where *T*<sup>r</sup> is the room temperature, and *T*<sup>m</sup> is the melt temperature of the material. *A*, *B*, *n*, *C*, and *m* are material–related parameters. The material parameters of S-7 tool steel and Q235 steel are presented in Table 5.

**Table 5.** Material constants for S-7 and Q235.


The orthotropic material model proposed by Long H. Nguyen et al. [14] was used for modeling the dynamic behavior of the UHMWPE layer subjected to ballistic impact. The material models consist of a nonlinear equation of the state of orthotropic, a strength model, and a failure model. The constitutive response of the material in the elastic regime is described as the orthotropic EOS composed of volumetric and deviatoric components. The pressure is defined by:

$$P = P(\varepsilon\_{\text{vol}}, e) - \frac{1}{3}(\mathsf{C}\_{11} + \mathsf{C}\_{21} + \mathsf{C}\_{31})\varepsilon\_{11}^{d} - \frac{1}{3}(\mathsf{C}\_{12} + \mathsf{C}\_{22} + \mathsf{C}\_{32})\varepsilon\_{22}^{d} \tag{5}$$

$$-\frac{1}{3}(\mathsf{C}\_{13} + \mathsf{C}\_{23} + \mathsf{C}\_{33})\varepsilon\_{33}^{d} \tag{6}$$

where *C*ij are the coefficients of the stiffness matrix, *ε d ij* refers to the deviatoric strains in the principal directions, and the volumetric component *P*(*εvol*,*e*) is defined by the Mie–Grüneisen EOS:

$$P(e\_{vol}, e) = P\_r(v) + \frac{\Gamma(v)}{v} [e - e\_r(v)] \tag{6}$$

where *v*, *e,* and Γ(*v*) represent the volume, internal energy, and the Grüneisen coefficient, respectively. *Pr*(*v*) is the reference pressure, and *er*(*v*) is the reference internal energy. The quadratic yield surface was adopted as the material strength model to describe the nonlinear, irreversible hardening behavior of the composite laminate:

$$f(\sigma\_{ij}) = a\_{11}\sigma\_{11}^2 + a\_{22}\sigma\_{22}^2 + a\_{33}\sigma\_{33}^2 + 2a\_{12}\sigma\_{11}\sigma\_{22} + 2a\_{23}\sigma\_{22}\sigma\_{33} \\ \tag{7}$$

$$+ 2a\_{13}\sigma\_{11}\sigma\_{33} + 2a\_{44}\sigma\_{23}^2 + 2a\_{55}\sigma\_{31}^2 + 2a\_{66}\sigma\_{12}^2 = k$$

where *a*ij are the plasticity coefficients, and *σ*ij represent the stresses in the principal directions of the material. In addition, the state variable, *k*, is used to define the border of the yield surface. It is described with a master and stress–effective plastic strain curve defined by ten piecewise points to consider the effect of strain hardening.

In the numerical models, the failure model of the orthotropic material is based on a combined stress criterion given as follows:

$$\left(\frac{\sigma\_{\text{ii}}}{S\_{\text{ii}}(1 - D\_{\text{ii}})}\right)^2 + \left(\frac{\sigma\_{\text{ii}}}{S\_{\text{ij}}(1 - D\_{\text{ii}})}\right)^2 + \left(\frac{\sigma\_{\text{ki}}}{S\_{\text{ki}}(1 - D\_{\text{ki}})}\right)^2 \ge 1 \text{ for } i, j, k = 1, 2, 3 \tag{8}$$

where *S* is the failure strength in the respective directions of the material, and *D* is the damage parameter following a linear relationship with stress and strain, as shown below:

$$D\_{\rm ii} = \frac{\mathcal{L}\sigma\_{\rm ii,f}\varepsilon\_{\rm cr}}{\mathcal{Z}\mathcal{G}\_{\rm ii,f}} \tag{9}$$

where *L* is the characteristic cell length, *ε*cr refers to the crack strain, and *G*ii,f presents the fracture energy in the direction of damage.

The constants for the orthotropic equation of state are presented in Table 6, and the parameters for orthotropic yield strength are shown in Table 7.


**Table 6.** Material constants for Orthotropic equation of state.

**Table 7.** Material constants for Orthotropic yield strength.


#### *4.2. Numerical Results and Analysis*

Table 8 presents the numerical simulation results of the A–P core penetrating the composite armor. *v*<sup>i</sup> and *v*<sup>r</sup> are the impacts and residual velocities of the ogive–nose penetrator. *p* is the depth of penetration. Due to the experimental results, the impact velocity is set from 430 m/s to 700 m/s. With the increased impact velocity, the penetration depth gradually increased. When the impact velocity reached 500 m/s, the ogive–nose penetrator could just perforate the composite armor.

**Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation. **Table 8.** Numerical simulation results of perforation.

The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the The contour of Von–Mises stress at the impact velocity of 500 m/s is shown in Figure 11. It can be inferred that the maximum stress exceeds the yield stress of the steel plate, and

Q235 back plate is pierced. Therefore, the velocity of 500 m/s can be considered as the ballistic limit of the composite armor, which is much higher than the 467.2 m/s obtained from the experimental results. The numerical simulation results are acceptable, with a relative error of 7.02%. ballistic limit of the composite armor, which is much higher than the 467.2 m/s obtained from the experimental results. The numerical simulation results are acceptable, with a relative error of 7.02%.

*Polymers* **2022**, *14*, x FOR PEER REVIEW 17 of 22

**Figure 11.** The contour of Von–Mises stress at the impact velocity of 500 m/s. **Figure 11.** The contour of Von–Mises stress at the impact velocity of 500 m/s.


The energy balance for the perforation is given by The energy balance for the perforation is given by

$$\frac{1}{2}mv\_{\text{i}}^{2} = \frac{1}{2}mv\_{\text{r}}^{2} + \mathcal{W} \tag{10}$$

$$\mathcal{W} = \mathcal{W}\_{\text{Q235}} + \mathcal{W}\_{\text{PE}} \tag{11}$$

where *m* is the mass of the projectile, *v*<sup>i</sup> is the impact velocity, *v*<sup>r</sup> is the residual velocity, and *W* is the work performed during perforation. The mass of the A–P core was set at 40.4 g, then the work conducted during the perforation of the composite armor could be calculated, as listed in Table 9. The value of *W* stayed stable from 5.05 kJ to 5.09 kJ, which means that dissipated energy in the petaling stays stable at around 5 kJ. At the ballistic limit from the numerical results, 500 m/s, the dissipated energy is the same as the work performed at a higher velocity after perforation. So, the principle of energy conservation can be applied here. where *m* is the mass of the projectile, *v*<sup>i</sup> is the impact velocity, *v*<sup>r</sup> is the residual velocity, and *W* is the work performed during perforation. The mass of the A–P core was set at 40.4 g, then the work conducted during the perforation of the composite armor could be calculated, as listed in Table 9. The value of *W* stayed stable from 5.05 kJ to 5.09 kJ, which means that dissipated energy in the petaling stays stable at around 5 kJ. At the ballistic limit from the numerical results, 500 m/s, the dissipated energy is the same as the work performed at a higher velocity after perforation. So, the principle of energy conservation can be applied here.

The Lambert–Jonas model [26,30–32] can provide a reasonable fit to predict the re-

sidual velocity of the penetrator after perforation. The model can be expressed as

**Table 9.** Results of calculated work *W* in the perforation. **Table 9.** Results of calculated work *W* in the perforation.

(2) Lambert–Jonas model


#### (2) Lambert–Jonas model *Polymers* **2022**, *14*, x FOR PEER REVIEW 18 of 22

The Lambert–Jonas model [26,30–32] can provide a reasonable fit to predict the residual velocity of the penetrator after perforation. The model can be expressed as

$$\begin{array}{l}\text{The performance of the measurement can be represented as} \\\\ v\_{\mathbf{r}} = \begin{cases} 0,0 \le v\_{\mathbf{i}} \le v\_{\mathbf{b}l} \\ a \left(v\_{\mathbf{i}}^{p} - v\_{\mathbf{b}l}^{p}\right)^{1/p}, v\_{\mathbf{i}} \ge v\_{\mathbf{b}l} \end{cases} \end{array} \tag{12}$$

where *v*<sup>i</sup> , *v*r, and *v*bl are the impact, residual, and ballistic limit velocity in normal impact. *α* and *p* are the coefficients, where 0 ≤ *α* ≤ 1 and *p* > 1. Based on the numerical simulation results, the Lambert–Jonas model can be established to predict the residual velocity of the A–P core after perforating the PE composite armor. where *v*i, *v*r, and *v*bl are the impact, residual, and ballistic limit velocity in normal impact. *α* and *p* are the coefficients, where 0 ≤ *α* ≤ 1 and *p* > 1. Based on the numerical simulation results, the Lambert–Jonas model can be established to predict the residual velocity of the A–P core after perforating the PE composite armor.

When the model with *p* = 2, the coefficient *α* can be set as 1, and the model can be justified based on the energy conservation law [33]. This model can be written as When the model with *p* = 2, the coefficient *α* can be set as 1, and the model can be justified based on the energy conservation law [33]. This model can be written as

$$v\_{\mathbf{r}} = \begin{cases} 0,0 \le v\_{\mathbf{i}} \le v\_{\mathbf{b}\mathbf{l}}\\ \left(v\_{\mathbf{i}}^2 - v\_{\mathbf{b}\mathbf{l}}^2\right)^{1/2}, v\_{\mathbf{i}} \ge v\_{\mathbf{b}\mathbf{l}} \end{cases} \tag{13}$$

the predicted *v*<sup>r</sup> − *v*<sup>s</sup> curve and the simulation results are presented below. As shown in Figure 12, the Lambert–Jonas model can be an effective method in predicting the residual velocity of the A–P core after perforation. In addition, the perforation process can be regarded as a rigid body penetration. the predicted *v*<sup>r</sup> − *v*<sup>s</sup> curve and the simulation results are presented below. As shown in Figure 12, the Lambert–Jonas model can be an effective method in predicting the residual velocity of the A–P core after perforation. In addition, the perforation process can be regarded as a rigid body penetration.

**Figure 12.** Comparison between the Lambert–Jonas model and the numerical simulations. **Figure 12.** Comparison between the Lambert–Jonas model and the numerical simulations.

#### (3) Cavity–Expansion Model (3) Cavity–Expansion Model

As the A–P core has a diameter of 12.48 mm and a length of 53.4 mm, the composite armor with a thickness of 53 mm can be considered an intermediate target. The square armor has a width of 300 mm, which is about 24 times the diameter of the A–P core. Thus, the cylindrical cavity expansion can be used to predict the ballistic limit of the A–P core. Figure 13 shows the dimensions of the A–P core. The caliber–radius–head (CRH) is 3.05, which is also denoted as *ψ*. As the A–P core has a diameter of 12.48 mm and a length of 53.4 mm, the composite armor with a thickness of 53 mm can be considered an intermediate target. The square armor has a width of 300 mm, which is about 24 times the diameter of the A–P core. Thus, the cylindrical cavity expansion can be used to predict the ballistic limit of the A–P core. Figure 13 shows the dimensions of the A–P core. The caliber–radius–head (CRH) is 3.05, which is also denoted as *ψ*.

**Figure 13.** Geometry and dimensions (in mm). **Figure 13.** Geometry and dimensions (in mm).

> Coefficient *k*<sup>1</sup> is expressed as Coefficient *k*<sup>1</sup> is expressed as

that  $k\_1$  is expressed as

$$k\_1 = \left(4\psi^2 - 4\psi/3 + 1/3\right) - \frac{4\psi^2(2\psi - 1)}{\sqrt{4\psi - 1}} \sin^{-1}\left[\frac{\sqrt{4\psi - 1}}{2\psi}\right] \tag{14}$$

The radial stress *σ*<sup>r</sup> at the cavity surface versus cavity expansion velocity *V* is given by [34] The radial stress *σ*<sup>r</sup> at the cavity surface versus cavity expansion velocity *V* is given by [34]

$$
\sigma\_r = \sigma\_s + \rho\_t \mathcal{B} V^2 \tag{15}
$$

 *r s t* = + *BV* (15) where *σ*<sup>s</sup> is the quasi–static radial stress required to open the cylindrical cavity, *ρ*<sup>t</sup> is the where *σ*<sup>s</sup> is the quasi–static radial stress required to open the cylindrical cavity, *ρ*<sup>t</sup> is the density of the target, and *B* is a dimensionless constant. *σ*s, *b*, and *B* are obtained from [23]

*b* = 1 − γ<sup>2</sup>

$$
\sigma\_{\mathbf{s}} = \frac{Y}{\sqrt{3}} \left\{ 1 + \left[ \frac{E}{\sqrt{3}Y} \right]^n \int\_0^b \frac{(-\ln x)^n}{1-\mathbf{x}} d\mathbf{x} \right\} \tag{16}
$$

$$
\mathbf{x}
$$

$$b = 1 - \gamma^2 \tag{17}$$

(17)

$$\nu = 1 \quad \text{y} \tag{17}$$

$$B = \frac{1}{2} \left\{ \frac{1}{(1-\nu)\sqrt{1-\alpha^2}} \ln \left[ \frac{1+\sqrt{1-\alpha^2}}{\alpha} \right] + \gamma^2 - 2\ln[\gamma] - 1 \right\} \tag{18}$$
 $\text{The yield stress and } \nu \text{ is Poisson's ratio of the target. } a \text{ and } \gamma \text{ are given by}$ 

 where *Y* is, the yield stress and *ν* is Poisson's ratio of the target. *α* and *γ* are given by

$$a^2 = \frac{\sqrt{3}(1 - 2\nu)}{2(1 - \nu)} \left(\frac{\rho\_l V^2}{Y}\right) \tag{19}$$
 
$$\Im(1 \pm \nu)Y$$

$$
\gamma^2 = \frac{2(1+\nu)Y}{\sqrt{3}E} \tag{20}
$$

<sup>2</sup> 2(1 ) 3 *Y E* + = (20) Furthermore, a rigid ogive–nosed projectile, with the impact velocity of *v*<sup>i</sup> , the ballistic limit of *v*bl and the residual velocity *v*r, is given by

$$v\_{bl} = \left(\frac{2\sigma\_s}{\rho\_p} \frac{h}{(L+k\_1 l)}\right)^{1/2} \left[1 + \mathcal{C} + \frac{2}{3}\mathcal{C}^2\right]^{1/2} \tag{21}$$

$$\left[\left(\frac{\rho\_h}{\rho\_h}\right)^2 - \left[\frac{1}{3}\frac{\rho\_h}{L}\right]^{1/2} \left[1 - \frac{1}{3}\mathcal{L}\right]^{1/2}\right]$$

$$v\_r = v\_{bl} \left[ \left( \frac{v\_i}{v\_{bl}} \right)^2 - 1 \right]^{1/2} \left[ 1 - \mathcal{C} + \frac{1}{2} \mathcal{C}^2 \right]^{1/2} \tag{22}$$

1/2 <sup>2</sup> 1/2 <sup>1</sup> <sup>2</sup> 1 1 *i r bl v v v C C* = − − + (22) where *C* is a small parameter related to the target inertia. When target inertia is neglected, the ballistic limit of *v*bl and the residual velocity *v*<sup>r</sup> can be simplified as [23,25,35] as

$$v\_{bl} = \left(\frac{2\sigma\_s}{\rho\_p} \frac{h}{(L+k\_1l)}\right)^{1/2} \tag{23}$$

$$v\_r = v\_{bl} \left[ \left( \frac{v\_i}{v\_{bl}} \right)^2 - 1 \right]^{1/2} \tag{24}$$

where the residual velocity *v*<sup>r</sup> is the same as the Lambert–Jones model in Equation (13).

Based on the constitutive models of the target materials, the quasi–static radial stress *σ*<sup>s</sup> can be expressed as [36]

$$
\sigma\_s = \frac{\chi}{\sqrt{3}} \left[ 1 + \ln \left( \frac{E}{\sqrt{3}\chi} \right) \right] + \frac{\pi^2 H}{18} \tag{25}
$$

where *E* and *H* are Young's modulus and the constant tangent modulus in the plastic region if the stress versus strain curve of the target can be expressed as

$$
\sigma = \begin{cases} \quad \text{E\varepsilon} \ \sigma < Y \\ Y + H \varepsilon \ \sigma \ge Y \end{cases} \tag{26}
$$

Thus, the value of *σ*<sup>s</sup> for the Q235 face sheets can be calculated. For UHMWPE laminates, there may not be a mature model to predict the quasi–static radial stress required to open the cylindrical cavity, but the range of the *σ*<sup>s</sup> can be estimated from the empirical formula [37,38] below,

$$
\sigma\_{\rm s} = (1.33 \sim 2) \mathbf{Y\_{\rm t}} \tag{27}
$$

When the coefficient is set as the minimum value of 1.33, the value at a relatively low level can be obtained, as listed in Table 10.


**Table 10.** The predicted value of quasi–static radial stress *σ*s.

For the composite armor composed of Q235 face sheets and UHMWPE laminates, the effective value of *σ*<sup>s</sup> can range from 2.76 GPa to 4.26 GPa. When the value of effective *σ*<sup>s</sup> is set as 3.08 GPa, the ballistic limit of the composite armor calculated from Equation (23) is 467 m/s, which is consistent with the value obtained from the numerical simulation results.

In conclusion, the principle of energy conservation and the Lambert–Jonas model can be applied to calculate the work performed during the perforation and the residual velocities of the A–P core after perforation. In addition, the quasi–static radial stress *σ*<sup>s</sup> required to open the cylindrical cavity can be estimated from the cavity–expansion model. With the value of 3.08 GPa, the predicted ballistic limit is consistent with the numerical simulation results.

#### **5. Conclusions**

A UHMWPE composite armor made up of two pieces of UHMWPE laminates in the middle and Q235 steel face sheets is proposed, and a study of the ballistic limit of the composite armor under the impact of a typical ogive–nose penetrator was carried out. (1) According to the experimental results, the average velocity of 501.1 m/s and 433.3 m/s can be taken as the ballistic limit of UHMWPE composite armor under the impact of the ogive–nose projectile, which is 467.2 m/s. In comparison, the ballistic limit obtained from the numerical simulation results was 500 m/s, which is acceptable with a relative error of 7.02%. (2) Petaling, as the main dominant failure mode for both face sheets, could be observed within the impact velocity range of 501 m/s to 1026 m/s. Delamination and shear failure dominated the penetration process of UHMWPE laminates. In addition, the charring layer on the front steel plate could be observed, and more severe ablation could

be noticed at the impact velocity of around 1000 m/s. (3) Through theoretical models, the perforation mechanism of composite armor under the impact of A–P cores was analyzed. The principle of energy conservation and the Lambert–Jonas model was applied to calculate the work performed during the perforation and the residual velocities. In addition, the quasi–static radial stress *σ*<sup>s</sup> required to open the cylindrical cavity were estimated from the cavity–expansion model. With the value of 3.08 GPa, the predicted ballistic limit was consistent with the numerical simulation results.

The ballistic limit of the UHMWPE composite armor under the impact of the ogive–nose projectile was considered to be 467.2 m/s, which indicates that the composite armor may not have a strong ability to resist the penetration of sharp head penetrators. In order to enhance the resistance against bullets such as A–P projectiles, UHMWPE should be strengthened, and the structure should be further optimized in future studies.

**Author Contributions:** Methodology, validation, formal analysis, writing—original draft preparation, funding acquisition, L.D.; supervision, visualization; writing—review, validation, X.G.; writing—review and editing, supervision, project administration, P.S.; supervision, data curation, resources, X.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** The authors thank the National Natural Science Foundation of China (Grant No. 11802142) and the project of the State Key Laboratory of Explosion Science and Technology (Grant No. KFJJ20-08M).

**Institutional Review Board Statement:** The study did not require ethical approval. This statement can be excluded.

**Data Availability Statement:** The raw and processed data generated during this study will be made available upon reasonable request.

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

#### **References**


## *Article* **Chemical and Enzymatic Fiber Modification to Enhance the Mechanical Properties of CMC Composite Films**

**Xiaobao Li 1,†, Zhengjie Tang 1,†, Zhenbing Sun <sup>1</sup> , John Simonsen <sup>2</sup> , Zhinan Luo <sup>1</sup> , Xiaoping Li 1,3,\* and Jeffery J. Morrell 4,\***


**Abstract:** Carboxymethyl cellulose (CMC) is a cellulose derivative that can be obtained from wood, bamboo, rattan, straw, and other cellulosic materials. CMC can be used to produce biofilms for many purposes, but the properties of these resulting films make them unsuitable for some applications. The effects of three kinds of plant fiber addition on CMC film properties was investigated using CMC derived from eucalyptus bark cellulose. Tensile strength (TS) and elongation at break (EB) of CMC/sodium alginate/glycerol composite films were 26.2 MPa and 7.35%, respectively. Tensile strength of CMC composite films substantially increased, reaching an optimum at 0.50 g of fiber. The enhancement due to industrial hemp hurd fiber on CMC composite films was more obvious. Pretreatment with hydrogen peroxide (H2O<sup>2</sup> ) and glacial acetic acid (CH3COOH) produced films with a TS of 35.9 MPa and an EB of 1.61%. TS values with pectinase pretreated fiber films was 41.3 MPa and EB was 1.76%. TS of films pretreated with pectinase and hemicellulase was 45.2 MPa and EB was 4.18%. Chemical and enzymatic treatment both improved fiber crystallinity, but film tensile strength was improved to a greater extent by enzymatic treatment. Surface roughness and pyrolysis residue of the film increased after fiber addition, but Fourier transform infrared spectroscopy (FTIR), opacity, and water vapor transmission coefficients were largely unchanged. Adding fiber improved tensile strength of CMC/sodium alginate/glycerol composite films and broadened the application range of CMC composite films without adversely affecting film performance.

**Keywords:** eucalyptus bark; Yunnan pine wood; bamboo culms; industrial hemp hurd; FTIR; XRD; TG; mechanical properties

## **1. Introduction**

Petroleum-based composite films are widely used in the food, pharmaceutical, and chemical industries due to their good properties and low cost. However, there is increasing interest in moving away from fossil fuel-based materials to renewable natural polymers such as cellulose. Cellulose can be modified to produce carboxymethyl cellulose (CMC), an odorless, tasteless, non-toxic, neutral or slightly alkaline, white or yellowish powder. CMC is hygroscopic, relatively light and heat stable, and transparent in aqueous solutions [1]. CMC is widely used in oil drilling, food packaging, concrete modification, and soil improvement [2–7]. The degree of substitution (DS, average number of hydroxyl groups substituted with carboxymethyl groups per anhydroglucose unit (AGU)) has a major influence on the properties, and therefore the potential uses, of CMC. The theoretical maximum DS of CMC is three; the degree of substitution of CMC directly affects the solubility, emulsification,

**Citation:** Li, X.; Tang, Z.; Sun, Z.; Simonsen, J.; Luo, Z.; Li, X.; Morrell, J.J. Chemical and Enzymatic Fiber Modification to Enhance the Mechanical Properties of CMC Composite Films. *Polymers* **2022**, *14*, 4127. https://doi.org/ 10.3390/polym14194127

Academic Editors: R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 17 August 2022 Accepted: 26 September 2022 Published: 2 October 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

thickening, stability, acid resistance, and salt resistance of CMC [8]. CMC with a super high degree of substitution (DS = 1.7~3.0) is often used in the textile, printing, and dyeing industry. CMC with a high degree of substitution (DS = 1.0~1.2) is often used as a food additive. CMC (DS = 0.6~0.9) with a low degree of substitution is commonly used in industrial drilling, ceramics, detergents, building materials, etc. [9]. In past research reports, some rich and underutilized plant cellulose sources have been used as raw materials for CMC production to replace cellulose materials obtained from cotton linter or wood bleached pulp. However, there is little information about CMC from eucalyptus bark cellulose. Eucalyptus bark comes from the eucalyptus tree, fallen off every year and is rich in cellulose. Eucalyptus wood is widely used in wood-based panel manufacturing, pulp and paper, etc., but the bark is a by-product of eucalyptus that typically has no use [10].

CMC also has great potential to create degradable films. However, single CMC composites are not suitable for all applications due to their poor mechanical properties and water permeability. Thickeners such as starch, sodium alginate, or gelatin can improve the mechanical properties and reduce moisture absorption of CMC films. When the ratio of CMC to corn starch is 4:6, the tensile strength of single starch film can be significantly improved, from 3.8 to 17.0 MPa [11], whereas addition of 1.5% sodium alginate to a CMC/chitosan mixture produced antibacterial food packaging films with tensile strengths and elongation at break of 65.32 MPa and 17.85%, respectively [12]. Adding 3.2% gelatin to a 0.8% CMC solution, the tensile strength of CMC composite film became 7.84 ± 0.30 MPa [13].Therefore, sodium alginate is often used as a thickener in CMC composite films. Plasticizers such as sorbitol, polyethylene glycol, or glycerin help improve ductility and tensile strength of composite samples. After 1.2% sorbitol was blended with CMC–gelatin–chitosan as plasticizer, the elongation at break of the composite film became 9.23% [14]. Adding 25% (*w*/*w*) glycerol to starch–alginate–CMC increased elongation at break by 58.6% [15]; 5 wt% polyethylene glycol blended with clay minerals–CMC, the maximum elongation at break was only 8.0% [16]. This comparison shows that the plasticizing effect of glycerin is the best.

CMC composite film properties can also be enhanced via addition of fibers derived from a variety of natural sources. For example, tensile strength increased 1.93 times with the addition of 8% sugar cane fiber to a polyvinyl alcohol (PVA) composite film [17]. The addition of wheat bran fiber to a corn starch composite film was associated with an increase in tensile strength from 2 to 5.07 MPa and a decrease in elongation at break from 60 to 28% [18]. The addition of cassava bagasse fiber to a cassava starch/glycerol composite film significantly increased maximum tensile strength (from 1.23 ± 0.15 to 7.78 ± 0.83 MPa) [19].The ability of small amounts of fiber to enhance film properties have seen these products used in construction, automotive, packaging, sports, and biomedicine. These applications highlight the potential for adding fibers to improve the properties of CMC films. However, there is little information about the effects of fiber-separation methods on the properties of CMC film. Fiber can be separated by mechanical or chemical methods. The mechanical processing method often results in ripped or torn fibers with a high elastic modulus and elongation at break, but poor tensile strength [20]. Chemical methods include nitric acid + potassium chlorate (HNO<sup>3</sup> + KClO3), sodium hypochlorite (NaClO), hydrogen peroxide + glacial acetic acid (H2O<sup>2</sup> + HAc), and sodium hydroxide (NaOH); sulfate can represent a gentler separation method [21]. For example, fibers prepared by H2O<sup>2</sup> + HAc were not hollow, and resulted in separated whole fibers with high tensile strength. There have been few reports on enzymatic separation of plant fibers; however, enzymes have catalytic efficiencies that are 107–10<sup>13</sup> times higher than nonenzymatic catalysts [22,23]. Enzymatic catalytic reactions are substrate-specific substrates, do not affect other raw materials, and cause less damage to raw materials. As enzymes are proteins, they are biodegradable, more environmentally friendly, and are a better choice for fiber separation [20,24,25].

In this study, chemical and enzymatic methods were used to separate plant fibers from different raw materials. The fibers were characterized and the properties of fiber-amended CMC composites were studied. We also suggest a method for making CMC films of high quality and improving the value of eucalyptus bark.

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

#### *2.1. Materials*

Eucalyptus bark (Eucalyptus globulus Labill.), Yunnan pine wood (Pinus yunnanensis), bamboo culms (Neosinocalamus affinis), and industrial hemp hurd (Cannabis sativa) were all obtained locally (Kunming, Yunnan Province, China). They were ground to pass through a 40–60 mesh screen (250–420 µm) before cellulose extraction and chemical analysis. The contents of benzene-alcohol extract, holo-cellulose, cellulose, hemicellulose, and lignin were determined according to Chinese Standards GB/T 2677.6-1994 (Determination of organic solvent extract in paper raw materials), GB/T 2677.10-1995 (Determination of holo-cellulose in paper raw materials), GB/T 744-1989 (Determination of α-cellulose in pulp), and GB/T 2677.8-94 (Determination of acid-insoluble lignin in paper raw materials).

### *2.2. Plant Fiber Separation*

Fibers of Yunnan pine wood, industrial hemp hurd, and bamboo culms were cut into pieces 3–5 by 3–5 by 5–10 mm (width by thickness by length) prior to chemical and enzymatic fiber extraction.

For chemical fiber separation: five grams of Yunnan pine wood, Bamboo culms, or industrial hemp hurd were placed in beakers and immersed in 100 mL of a 50:50 mixture of hydrogen peroxide (H2O2): glacial acetic acid (CH3COOH) at 70 ◦C until the sample turned white. The samples were washed with distilled water until the pH was 7, then the samples were shaken slightly to separate the fibers. These procedures were performed in triplicate.

Enzymatic fiber separation: three enzymatic methods were used to obtain fibers. In the first method, five grams of a given material was treated with 100 mL of 5.00% lipase solution at 50 ◦C, stirred (900–1000 r/min) for 3 h, then the solution was filtered off and the residual materials were treated with GB2677.10-1995 (Determination of the content of holo-cellulose from paper raw materials) to remove most of the lignin. The residual material was immersed in 40 mL of 50 ◦C distilled water and stirred (900–1000 r/min) for more than 3 h until most of the fibers were separated, then washed with distilled water to obtain residues. The second method used the same procedure, but then immersed the fibers in 40 mL of 5% pectinase solution instead of distilled water. The final method used the same procedure, but then immersed the materials in 40 mL of 5% a 20:20 mixture of 5% pectinase solution and 5% hemicellulase solution instead of distilled water. Lipase (CAS:9001-62-1), pectinase (CAS:9032-75-1), and hemicellulose (CAS:9025-56-3) were obtained from Aladdin Biotechnology Co., Ltd., (Shanghai, China) and had enzyme activities of 100,000; 30,000; and 5000 U/g, respectively. Each fiber material/enzymatic treatment combination was prepared in triplicate.

#### *2.3. Preparation of Cellulose, CMC, and CMC Composite Films*

Four grams of cellulose extracted from eucalyptus bark according to Chinese Standard GB/T744-1989 (Determination of α-cellulose from pulp) was mixed with 80 mL of 100% ethanol(CH3CH2OH) and 20 mL of 30% NaOH solution and then stirred(900–1000 r/min) for 60 min at 30 ◦C. Five grams of sodium chloroacetate (C2H2ClNaO2) were added and heated at 65 ◦C for 3 h. The sample was washed with 90% glacial acetic acid to a pH of 7, then washed with 80% ethanol 3 times and 95% ethanol once, before being oven-dried at 65 ◦C for 3 h to obtain CMC.

One gram of CMC was mixed with 0.40 g sodium alginate (C6H7NaO6), 0.15 g glycerol (C3H8O3), and 49.00 g distilled water and stirred (900–1000 r/min) at 70 ◦C for 15 min. The mixture was treated in an ultrasonic bath for 10 min and vacuumed for 45 min to remove air bubbles. The solution was poured into a polytetrafluoroethylene (PTFE) mold and cured at 30 ◦C for 48 h to obtain CMC composite films. The fiber-modified CMC composite film

*Polymers* **2022**, *14*, x FOR PEER REVIEW 4 of 15

process was similar except that 0.1, 0.3, and 0.5 g of a given plant fiber was added during mixing (shown in Figure 1). posite film process was similar except that 0.1, 0.3, and 0.5 g of a given plant fiber was added during mixing (shown in Figure 1).

One gram of CMC was mixed with 0.40 g sodium alginate (C6H7NaO6), 0.15 g glycerol (C3H8O3), and 49.00 g distilled water and stirred (900–1000 r/min) at 70 °C for 15 min. The mixture was treated in an ultrasonic bath for 10 min and vacuumed for 45 min to remove air bubbles. The solution was poured into a polytetrafluoroethylene (PTFE) mold and cured at 30 °C for 48 h to obtain CMC composite films. The fiber-modified CMC com-

**Figure 1.** Preparation of CMC composite films. **Figure 1.** Preparation of CMC composite films.

#### *2.4. Fiber and Composite Film Characterization 2.4. Fiber and Composite Film Characterization*

**Fiber dimensions:** The length and width of 100 fibers, as well as the cell wall thickness and cell cavity width of each of the chemically or enzymatically prepared fibers were measured under a light microscope by ImageJ software. **Fiber dimensions:** The length and width of 100 fibers, as well as the cell wall thickness and cell cavity width of each of the chemically or enzymatically prepared fibers were measured under a light microscope by ImageJ software.

**Sample microstructure:** The microstructure of eucalyptus bark powder, cellulose, CMC, and CMC composite films were observed by placing a sample on an aluminium stub and coating with gold/palladium before observation with the Czech TESCAN MIRA LMS field emission scanning electron microscope at an accelerating voltage of 200 eV to **Sample microstructure:** The microstructure of eucalyptus bark powder, cellulose, CMC, and CMC composite films were observed by placing a sample on an aluminium stub and coating with gold/palladium before observation with the Czech TESCAN MIRA LMS field emission scanning electron microscope at an accelerating voltage of 200 eV to 30 KeV. A minimum of five fields were examined per material.

30 KeV. A minimum of five fields were examined per material. **Fourier-transform infrared spectroscopy (FTIR):** Eucalyptus bark powder, cellulose, and CMC were mixed with KBr and formed into a pellet while the CMC composite films were directly analyzed on a Nicolet i50 FTIR analyzer (Thermo Nicolet Corporation, Madison, WI, USA) with a scanning range of 500 to 4000 cm−1 and 64 scans. Baseline correction **Fourier-transform infrared spectroscopy (FTIR):** Eucalyptus bark powder, cellulose, and CMC were mixed with KBr and formed into a pellet while the CMC composite films were directly analyzed on a Nicolet i50 FTIR analyzer (Thermo Nicolet Corporation, Madison, WI, USA) with a scanning range of 500 to 4000 cm−<sup>1</sup> and 64 scans. Baseline correction was performed to analyze the spectral differences between plant fibers obtained by different treatments.

was performed to analyze the spectral differences between plant fibers obtained by different treatments. **X-ray diffractometer analysis (XRD):** The crystal structures of eucalyptus bark, cellulose, and the CMC films were studied by X-ray diffraction (XRD) on an Ultima IV X-ray **X-ray diffractometer analysis (XRD):** The crystal structures of eucalyptus bark, cellulose, and the CMC films were studied by X-ray diffraction (XRD) on an Ultima IV X-ray diffractometer (Rigaku Corporation, Tokyo, Japan) using a scanning angle from 5 to 60◦ , a step size of 0.026◦ (accelerating current = 30 mA and voltage = 40 kV), and Cu-Kα radi-

diffractometer (Rigaku Corporation, Tokyo, Japan) using a scanning angle from 5 to 60°,

ation of *λ* = 0.154 nm. The degree of crystallinity (DOC, %) was calculated according to the formula:

$$\text{DOC}\% = \frac{I\_{\text{Max}} - I\_{Am}}{I\_{\text{Max}}} \times 100 \tag{1}$$

*IMax* is the maximum intensity of the main peak (about 22◦ ), and *IAm* is the diffraction intensity of amorphous cellulose (about 15◦ ).

**Thermogravimetric (TG) analysis:** Approximately 5 to 6 mg of sample powder ground to pass through an 80 to 120 mesh screen and placed into sample holders for analysis on a TGA92 thermo gravimetric analyzer (KEP Technologies EMEA, Caluire, France). N<sup>2</sup> was used as the shielding gas and Al2O<sup>3</sup> as the reference compound. The temperature was increased from 35 to 800 ◦C at a rate of 20 ◦C/min to generate a thermogravimetric curve.

**Degree substitution (DS) of CMC:** The degree of substitution of hydroxyl groups has an important influence on resulting CMC properties. The degree of substitution was determined by the acidimeter method by weighing 0.2 g (accuracy 0.1 mg) of the sample, dissolving it in 80 mL of water, stirring for 10 min, and adjusting the pH to 8.0. The sample was titrated using sulphuric acid (H2SO4) with continuous stirring to pH 3.74. The volume (mL) of sulphuric acid titration solution used was recorded (to the nearest 0.05 mL). The degree of substitution (*DS*) was then calculated using the amount required to reach the end point according to Equations (1) and (2), as follows (Table 1).

$$B = \frac{2cV}{m} \tag{2}$$

$$DS = \frac{0.162B}{1 - 0.08B} \tag{3}$$

where *B* = amount of carboxymethyl substance contained in the sample, mmol/g; *m* = quality of the sample, g;

*c* = concentration of sulphuric acid standard titration solution, mol/L;

*V*= volume value of standard titration solution of sulphuric acid, mL.

**Table 1.** Chemical composition of eucalyptus bark, Yunnan pine wood, bamboo culms, and industrial hemp hurd <sup>a</sup> .


<sup>a</sup> Values represent means of 3 replicates. The numbers in parentheses are one standard deviations.

**Physical Properties:** Tensile strength (MPa) and elongation at break (%) were measured on ten 0.089 to 0.098 mm by 150 mm dog-bone samples of each material on a universal testing machine according to procedures described in GB/T 1040.1-2006 (Plastics— Determination of tensile properties). The load was applied to failure at a rate of 1 mm/min.

**Film Opacity and Water Vapor Transmission:** Opacity of the CMC composite films was tested by cutting 10 by 40 mm samples and placing them on the inner surface on one side of a cuvette and measuring absorbance at 600 nm on an XP Spectrum 752# ultraviolet spectrophotometer (XP-Spectrum Company, Shanghai, China). Five measurements were made for each sample.

Water vapor transmission rate was assessed under controlled temperature and relative humidity conditions using unit time, unit water vapor pressure difference, and thickness through the unit, and expressed as the unit area of the water vapor volume of the specimen.

The water vapor transmission coefficient of the specimen was calculated according to Equation (3).

$$\mathbf{P} = \frac{\Delta \mathbf{m} \times \mathbf{d}}{\mathbf{A} \times \mathbf{t} \times \Delta \mathbf{p}} \tag{4}$$

where P is the water vapor transmission coefficient of the sample in grams/square centimeter per second Pascal [g cm/(cm<sup>2</sup> ·s·Pa)].

∆m is the mass change of the sample in grams (g) during the period t.

A is the sample area through the water vapor in square meters (m<sup>2</sup> ).

t is the difference in time between two intervals after the mass change has stabilized in hours (h).

d is the thickness of the specimen in centimeters (cm).

∆p is the difference in water vapor pressure between the two sides of the specimen in Pascals (Pa).

#### *2.5. Statistical Analysis*

Equality of variance was confirmed using Fisher's test for raw material chemical analyses, fiber size measurements, and physical properties measurements of CMC composite films. Student's t test was carried out to compare the samples in pairs at *p* < 0.05. Data were analyzed using SPSS 25.0 statistical package (IBM, Armonk, NY, USA).

#### **3. Results and Discussion**

#### *3.1. Chemical Composition*

Cellulose content of eucalyptus bark was 44.9% higher than that of Yunnan pine wood and lower than that of bamboo culms or industrial hemp hurd (Table 1). Hemicellulose content of eucalyptus bark was 26.6% and was higher than that of other three materials. Lignin content of the bark was 27.2% lower than that of Yunnan pine wood and higher than either bamboo culms or industrial hemp hurd. These results were consistent with previous research and indicate that eucalyptus bark-derived cellulose is a suitable alternative [26]. Using this material as a substitute would reduce chemical consumption and allow utilization of a waste product.

#### *3.2. Effects of Different Pretreatment Methods on Fiber Yield and Dimensions*

Chemical treatment resulted in 100% fiber yield after 10 h of treatment at 70 ◦C (Table 2). Enzyme treatments produced lower yields. Distilled water treatment resulted in less than 5% fiber yields from bamboo culms and industrial hemp hurd and only 10–20% yield from Yunnan pine wood, despite the 12 h total treatment time. The subsequent use of pectinase alone or in combination with hemicellulase resulted in 90 to 95% fiber yield. Pectin plays important roles in cell wall interactions, especially in primary cell wall formation, and its disruption may facilitate fiber separation. Treatment times for the pectinase treatments were only 6 to 8 h whereas they were 5.5 to 8.5 h for the pectinase and hemicellulase treatments. These results illustrate the potential for producing high fiber yields using enzymatic treatments [27].

The distilled water treatment resulted in little fiber recovery and will not be further discussed. The other treatments had varying effects on the properties of the resulting fibers (Table 3). Pectinase treatment resulted in the shortest Yunnan pine wood fibers whereas the chemical treatment resulted in fibers that were nearly 50% longer. Similarly, pectinase and hemicellulase treated bamboo culms and industrial hemp hurd fibers were only half as long as those from the chemical treatment. Fiber length tended to be greater in all of the chemical treatments compared with the enzymatic treatments, although the pectinase treatment was sometimes similar to the chemical treatment (Table 3). The largest difference in fiber length was found between chemically and enzymatically treated industrial hemp hurd fibers. Decreased fiber lengths may reflect a tendency for enzymatically treated fibers to break more easily as they are separated, which would be detrimental to increasing the tensile strength of any composite. Fiber widths and lumen diameters tended to be similar

for the same material regardless of whether the samples were chemically or enzymatically treated. The treatments are less likely to affect fiber width or lumen size, given that their primary effects would be on the cell walls themselves. Cell walls were slightly thicker in pectinase-treated Yunnan pine wood fibers than chemically treated fibers, whereas cell wall thicknesses of pectinase and hemicellulase and chemically treated Yunnan pine wood fibers were similar. Pectin is an important component in primary wood cell wall formation but becomes less important with subsequent lignification. However, the specificity of pectinase for pectin could lead to more efficient separation with reduced breakage.

**Table 2.** Effect of chemical treatment alone or coupled with three sequential enzyme treatments on fiber yield and total treatment time of Yunnan pine wood, bamboo culms, and industrial hemp hurd fibers.


<sup>a</sup> Values represent results from three replicates per material per treatment.

**Table 3.** Effect of chemical or enzymatic treatment on characteristics of fibers derived from Yunnan pine wood, bamboo culms, and industrial hemp hurd <sup>a</sup> .


<sup>a</sup> Values represent means of 100 fibers per material per treatment while figures in parentheses are the range.

Fiber length to width ratios can be a useful indicator of potential effects of fiber addition on tensile properties. The addition of fibers with higher length to width ratios may have a greater effect on tensile strength. Length to width ratios tended to be smaller in enzymatically treated fibers than in chemically treated fibers of the same species. As noted earlier, this may reflect a tendency for enzymatically treated fibers to be more brittle and produce shorter fibers, which would reduce tensile properties.

The relative crystallinity of untreated Yunnan pine wood, bamboo culms, and industrial hemp hurd were 38.8, 49.1, and 47.4%, respectively (Table 4). Almost all of the lignin was removed from materials treated chemically or enzymatically and crystallinity was increased. The degree of crystallinity was greatest in Yunnan pine wood but crystallinity also increased in bamboo culms and industrial hemp hurd, although the differences were not significant. Increased crystallinity indicates that chemical and enzymatic treatments

removed some amorphous cellulose, resulting in an increase in the proportion of crystalline cellulose. Increased crystallinity may result in stronger reinforcing fibers.

**Table 4.** Effect of chemical or enzymatic treatment on relative crystallinity of Yunnan pine wood, bamboo culms, and industrial hemp hurd fibers <sup>a</sup> .


<sup>a</sup> Values represent means of 3 replicates.

#### *3.3. CMC Characterization*

**Degree of Substitution:** The substitution degree (DS) on CMC from eucalyptus bark cellulose was 0.89, which is similar to the values obtained for corn stover, straw, and reed CMCs, which ranged from 0.6 to 1.0 [28,29]. *3.3. CMC Characterization Degree of Substitution:* The substitution degree (DS) on CMC from eucalyptus bark cellulose was 0.89, which is similar to the values obtained for corn stover, straw, and reed CMCs, which ranged from 0.6 to 1.0 [28,29].

**SEM:** SEM examination of eucalyptus bark revealed that it consisted of many substances tightly aggregated together in small granular form (Figure 2a-1). The cellulose recovered from this material was in the form of polymerized fibrous bundles (Figure 2a-2), which became more discrete when the materials were reacted to form CMC with differing degrees of fiber breakage (Figure 2a-3). These changes reflect the effects of alkaline treatment and subsequent esterification, coupled with water penetration into the cellulose bundles, with resulting chain separation. *SEM:* SEM examination of eucalyptus bark revealed that it consisted of many substances tightly aggregated together in small granular form (Figure 2a-1). The cellulose recovered from this material was in the form of polymerized fibrous bundles (Figure 2a-2), which became more discrete when the materials were reacted to form CMC with differing degrees of fiber breakage (Figure 2a-3). These changes reflect the effects of alkaline treatment and subsequent esterification, coupled with water penetration into the cellulose bundles, with resulting chain separation.

**Figure 2.** (**a-1**) SEM images showing microstructure differences for unprocessed eucalyptus bark (**a-2**), cellulose microfibrils after chemical treatment, and (**a-3**) CMC produced from the eucalyptus bark cellulose. **Figure 2.** (**a-1**) SEM images showing microstructure differences for unprocessed eucalyptus bark (**a-2**), cellulose microfibrils after chemical treatment, and (**a-3**) CMC produced from the eucalyptus bark cellulose.

*FTIR Analysis:* FTIR spectra of commercially available CMC and eucalyptus bark CMC both contained the stretching vibrations of the CMC carboxylate anion COO at about 1630 cm−1, with two characteristic absorption peaks at 1410 cm−1 corresponding to the asymmetric (C=O) and symmetric (C-O) stretching vibrations caused by the carboxylic acid group, and an absorption peak at 1030 cm−1 corresponding to the stretching vibration of the cellulose C-O-C group. The absorption peak at 898 cm−1 is characteristic of the βglycosidic bond in cellulose [30–37]. No absorption peaks were observed at 1518 and 1320 cm−1, which are attributed to the aromatic vibration of the lignin ring and the C-O stretching vibration of the syringyl group, respectively [38], nor was there a peak at 1730 cm−1, which is attributed to hemicellulose. These results indicate that cellulose extraction was nearly complete with little evidence of residual lignin or hemicelluloses, and the cellulose was successfully transformed into CMC (Figure 3). **FTIR Analysis:** FTIR spectra of commercially available CMC and eucalyptus bark CMC both contained the stretching vibrations of the CMC carboxylate anion COO- at about 1630 cm−<sup>1</sup> , with two characteristic absorption peaks at 1410 cm−<sup>1</sup> corresponding to the asymmetric (C=O) and symmetric (C-O) stretching vibrations caused by the carboxylic acid group, and an absorption peak at 1030 cm−<sup>1</sup> corresponding to the stretching vibration of the cellulose C-O-C group. The absorption peak at 898 cm−<sup>1</sup> is characteristic of the β-glycosidic bond in cellulose [30–37]. No absorption peaks were observed at 1518 and 1320 cm−<sup>1</sup> , which are attributed to the aromatic vibration of the lignin ring and the C-O stretching vibration of the syringyl group, respectively [38], nor was there a peak at 1730 cm−<sup>1</sup> , which is attributed to hemicellulose. These results indicate that cellulose extraction was nearly complete with little evidence of residual lignin or hemicelluloses, and the cellulose was successfully transformed into CMC (Figure 3).

**Figure 3.** FTIR spectra of bark, holo−cellulose, cellulose, and CMC prepared from eucalyptus bark, as well as commercially available CMC. **Figure 3.** FTIR spectra of bark, holo−cellulose, cellulose, and CMC prepared from eucalyptus bark, as well as commercially available CMC. **Figure 3.** FTIR spectra of bark, holo−cellulose, cellulose, and CMC prepared from eucalyptus bark, as well as commercially available CMC.

*X-ray Diffraction*: Eucalyptus bark and holo-cellulose showed diffraction peaks at 2θ = 15.65, 16.45, 22.15, and 34.35°, whereas the cellulose diffraction peaks were at 14.85, 15.4, 21.75, and 34.35°, respectively (Figure 4). All three of the latter peaks correspond to the crystal planes of (101), (101ത), (200), and (004), which are typical reflections of cellulose Type I [39–43]. The relative crystallinity of eucalyptus bark and cellulose were 46.39 and 55.07%, respectively, again indicating that the cellulose was successfully extracted. The diffraction peaks of bark-derived CMC were at 2θ = 19.95 and 31.8°, which were similar to the diffraction peaks obtained from commercially available CMC, indicating successful transformation of cellulose to CMC (Figure 4) [4,43]. **X-ray Diffraction**: Eucalyptus bark and holo-cellulose showed diffraction peaks at 2θ = 15.65, 16.45, 22.15, and 34.35◦ , whereas the cellulose diffraction peaks were at 14.85, 15.4, 21.75, and 34.35◦ , respectively (Figure 4). All three of the latter peaks correspond to the crystal planes of (101), (101), (200), and (004), which are typical reflections of cellulose Type I [39–43]. The relative crystallinity of eucalyptus bark and cellulose were 46.39 and 55.07%, respectively, again indicating that the cellulose was successfully extracted. The diffraction peaks of bark-derived CMC were at 2θ = 19.95 and 31.8◦ , which were similar to the diffraction peaks obtained from commercially available CMC, indicating successful transformation of cellulose to CMC (Figure 4) [4,43]. *X-ray Diffraction*: Eucalyptus bark and holo-cellulose showed diffraction peaks at 2θ = 15.65, 16.45, 22.15, and 34.35°, whereas the cellulose diffraction peaks were at 14.85, 15.4, 21.75, and 34.35°, respectively (Figure 4). All three of the latter peaks correspond to the crystal planes of (101), (101ത), (200), and (004), which are typical reflections of cellulose Type I [39–43]. The relative crystallinity of eucalyptus bark and cellulose were 46.39 and 55.07%, respectively, again indicating that the cellulose was successfully extracted. The diffraction peaks of bark-derived CMC were at 2θ = 19.95 and 31.8°, which were similar to the diffraction peaks obtained from commercially available CMC, indicating successful transformation of cellulose to CMC (Figure 4) [4,43].

**Figure 4.** XRD diffractogram showing bark, holo-cellulose, cellulose, and CMC from eucalyptus bark, as well as commercially available CMC. **Figure 4.** XRD diffractogram showing bark, holo-cellulose, cellulose, and CMC from eucalyptus bark, as well as commercially available CMC. **Figure 4.** XRD diffractogram showing bark, holo-cellulose, cellulose, and CMC from eucalyptus bark, as well as commercially available CMC.

## *3.4. Effects of Fiber Treatment Method and Addition of Composite Film Properties*

The addition of chemically treated Yunnan pine wood or bamboo culm fibers to the CMC/sodium alginate/glycerol film was associated with decreased tensile strength with increasing fiber content, although the difference in bamboo culm was small. Conversely, tensile strength of the CMC composite film increased with increased industrial hemp hurd fiber. The addition of enzyme-treated Yunnan pine wood, bamboo culms, or industrial hemp hurd fibers enhanced tensile strength of CMC composite films, although this enhancement was small at the lowest addition level (0.1% wt/wt). The results showed that enzymatically treated plant fibers were more likely to produce better CMC composite films (Table 5). Industrial hemp hurd fiber has a smaller wall lumen ratio than Yunnan pine and bamboo fibers. Small compounds such as CMC, sodium alginate, or glycerol are more likely to enter the fiber lumen, especially on enzymatically modified shorter fibers. This should improve tensile strength of the resulting composite.

**Table 5.** Effect of addition of chemically or enzymatically derived Yunnan pine wood, bamboo culms, and industrial hemp hurd fibers to CMC film on the tensile strength, elongation at break, opacity, or vapor transmission <sup>a</sup> .


<sup>a</sup> Values represent means of replicates per treatment and figures in parentheses represent one standard deviation. Values followed by the same letter(s) do not differ-significantly from one another by Tukey's pairwise comparisons (α = 0.05).

The addition of chemically or enzymatically treated fibers produced more variable results with elongation at break. Values dropped sharply from the non-modified control but there were no consistent trends associated with fiber additive level or chemical vs. enzymatic pre-treatment. Although fiber addition clearly altered elongation at break, there were no consistent trends with regard to pre-treatment method or additive level (Table 5). Plant fibers usually have high rigidity and should improve tensile strength, but elongation at break could also decrease sharply, as observed in other materials [44].

Opacity is a useful measure for assessing the suitability of a film for commercial purposes. Although opacity varied widely with fiber pre-treatment and additive level, there were no consistent trends with regard to pre-treatment method or concentration. These results suggest that addition of low levels of chemically or enzymatically treated fibers had no consistent effect on opacity, likely because the overall fiber levels remained low.

The background water vapor coefficient for non-modified CMC was 0.20 g·cm/(cm<sup>2</sup> ·s·Pa), which is in line with previous reports for CMC films [45]. Addition of 0.1, 0.3, or 0.5 g of chemically or enzymatically recovered fibers to the CMC film had no significant effect on water vapor transmission regardless of plant source (Table 5). These results suggest that these fibers have the potential to improve tensile strength without negatively impacting the ability of the film to function as a water barrier. These attributes would make the films more suitable for food storage.

The data were subjected to an analysis of variance and means were examined using a Tukey's pairwise comparison test (α = 0.05).

**Pyrolytic Properties**: The addition of Yunnan pine wood, bamboo culms, or industrial hemp hurd obtained by pectinase and hemicellulase treatment to the CMC was associated with a higher first decomposition peak compared with the non-modified CMC (Table 6). Mass loss at that time was lower for the Yunnan pine wood and bamboo culms but similar for the industrial hemp hurd. The second decomposition peak was slightly lower with the addition of either Yunnan pine wood or bamboo culms fiber but nearly the same as non-amended CMC with addition of industrial hemp hurd. The mass losses at the second decomposition peak were all slightly higher with addition of Yunnan pine wood and industrial hemp hurd losing the most mass. The addition of fibers to the CMC altered both peak temperatures, whereas it decreased mass losses for the first peak and increased them for the second. The final residual weights of the films were 28, 32, 35, and 22 wt%, respectively. The addition of plant fibers resulted in an increase in the required decomposition temperature and an increase in the final weight residue (Figure 5) [4,46].

**Table 6.** Effect of addition of Yunnan pine wood, bamboo culms, and industrial hemp hurd fibers obtained using pectinase + hemicellulase treatment on decomposition peaks and mass losses of CMC film.


**Figure 5.** Effect of addition of Yunnan pine wood, bamboo culms, or industrial hemp hurd fibers obtained by pectinase + hemicellulase treatment on (**A**) TG and (**B**) DTG curves of eucalyptus bark CMC composite films. **Figure 5.** Effect of addition of Yunnan pine wood, bamboo culms, or industrial hemp hurd fibers obtained by pectinase + hemicellulase treatment on (**A**) TG and (**B**) DTG curves of eucalyptus bark CMC composite films.

#### **4. Conclusions**

(D21027).

**4. Conclusions**  Fibers prepared by chemical and enzymatic methods differed in length, width, cell wall thickness, and lumen width, which improved the crystallinity of the fibers. The length:width ratio of the enzymatically prepared fibers was smaller than that of chemically prepared fibers. Wall thickness:lumen width ratios of industrial hemp hurd fibers were the smallest among the three materials. Tensile strength (TS) and elongation at break (EB) of CMC composite films without plant fiber were 26.2 MPa and 7.35%, respectively. TS of CMC composite films was greatly improved by addition of industrial hemp hurd fiber, especially after enzyme treatment. TS of CMC composite films increased to 35.9, 41.3, and 45.2 MPa for chemical, pectinase, and pectinase + hemicellulase treatments, respectively, after adding 0.5 g industrial hemp hurd fiber. EB changed to 1.61, 1.76, and 4.18%, respectively. Addition of modified fibers did not affect opacity or water vapor per-Fibers prepared by chemical and enzymatic methods differed in length, width, cell wall thickness, and lumen width, which improved the crystallinity of the fibers. The length:width ratio of the enzymatically prepared fibers was smaller than that of chemically prepared fibers. Wall thickness:lumen width ratios of industrial hemp hurd fibers were the smallest among the three materials. Tensile strength (TS) and elongation at break (EB) of CMC composite films without plant fiber were 26.2 MPa and 7.35%, respectively. TS of CMC composite films was greatly improved by addition of industrial hemp hurd fiber, especially after enzyme treatment. TS of CMC composite films increased to 35.9, 41.3, and 45.2 MPa for chemical, pectinase, and pectinase + hemicellulase treatments, respectively, after adding 0.5 g industrial hemp hurd fiber. EB changed to 1.61, 1.76, and 4.18%, respectively. Addition of modified fibers did not affect opacity or water vapor permeability, indicating that adding low levels of fiber to CMC significantly improved film characteristics making them potentially suitable for food storage.

characteristics making them potentially suitable for food storage. **Author Contributions:** Conceptualization, X.L. (Xiaoping Li) and J.J.M.; methodology, X.L. (Xiaoping Li) and J.J.M.; validation, X.L (Xiaobao Li). Z.T., Z.S., Z.L. and J.S.; formal analysis, X.L. (Xiaoping Li) and J.J.M.; investigation, X.L. (Xiaobao Li), Z.T., Z.S., Z.L. and J.S.; resources, X.L. (Xiaoping Li); data curation, Z.T.; writing—original draft preparation, X.L. **Author Contributions:** Conceptualization, X.L. (Xiaoping Li) and J.J.M.; methodology, X.L. (Xiaoping Li) and J.J.M.; validation, X.L. (Xiaobao Li). Z.T., Z.S., Z.L. and J.S.; formal analysis, X.L. (Xiaoping Li) and J.J.M.; investigation, X.L. (Xiaobao Li), Z.T., Z.S., Z.L. and J.S.; resources, X.L. (Xiaoping Li); data curation, Z.T.; writing—original draft preparation, X.L. (Xiaobao Li); writing—review and editing, J.J.M., X.L. (Xiaoping Li); supervision, X.L. (Xiaoping Li); project administration, X.L. (Xiaoping Li) and J.J.M.; funding acquisition, X.L. (Xiaoping Li). All authors have read and agreed to the published version of the manuscript.

meability, indicating that adding low levels of fiber to CMC significantly improved film

(Xiaobao Li); writing—review and editing, J.J.M., X.L. (Xiaoping Li); supervision, X.L. (Xiaoping Li); project administration, X.L. (Xiaoping Li) and J.J.M.; funding acquisition, X.L. (Xiaoping Li). All authors have read and agreed to the published version of the man-**Funding:** This study was supported by the National Nature Science Foundation (31870551), Top Young Talents in Yunnan Province (YNWR-QNBJ-2018-120) and 111 Project (D21027).

uscript. **Institutional Review Board Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available from the listed authors.

**Funding:** This study was supported by the National Nature Science Foundation (31870551), Top Young Talents in Yunnan Province (YNWR-QNBJ-2018-120) and 111 Project **Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


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


## *Article* **Effects of Polypropylene Fibers on the Frost Resistance of Natural Sand Concrete and Machine-Made Sand Concrete**

**Yan Tan 1,\*, Junyu Long <sup>1</sup> , Wei Xiong <sup>2</sup> , Xingxiang Chen <sup>1</sup> and Ben Zhao <sup>1</sup>**


**Abstract:** In order to study the effect of polypropylene fibers on the frost resistance of natural sand and machine-made sand concrete, polypropylene fibers (PPF) of different volumes and lengths were mixed into natural sand and machine-made sand concrete, respectively. The freeze–thaw cycle test was carried out on polypropylene-fiber-impregnated natural sand concrete (PFNSC) and polypropylene-fiber-impregnated manufactured sand concrete (PFMSC), respectively, and the apparent structural changes before and after freezing and thawing were observed. Its strength damage was analyzed. A freeze–thaw damage model and a response surface model (RSM) were established used to analyze the antifreeze performance of PFMSC, and the effects of the fiber content, fiber length, and freeze–thaw times on the antifreeze performance of PFMSC were studied. The results show that with the increase in the number of freeze–thaw cycles, the apparent structures of the PFMSC gradually deteriorated, the strength decreased, and the degree of freeze–thaw damage increased. According to the strength damage model, the optimum volume of PPF for the PFNSC specimens is 1.2%, and the optimum volume of PPF for the PFMSC specimens is 1.0%. According to the prediction of RSM, PFNSC can maintain good antifreeze performance within 105 freeze–thaw cycles, and when the PPF length is 11.8 mm, the antifreeze performance of PFNSC reaches the maximum, its maximum compressive strength value is 33.8 MPa, and the split tensile strength value is 3.1 MPa; PFMSC can maintain a good antifreeze performance within 96 freeze–thaw cycles. When the length of PPF is 9.1 mm, the antifreeze performance of PFMSC reaches the maximum, its maximum compressive strength value is 45.8 MPa, and its split tensile strength value is 3.2 MPa. The predicted values are in good agreement with the measured values, and the model has high reliability.

**Keywords:** polypropylene fiber mechanism sand concrete; frost resistance; freeze–thaw damage model; RSM strength model

## **1. Introduction**

Durability is an important indicator for measuring the ability of a material to resist the long-term damage of both itself and the natural environment [1–4]. Freeze– thaw damage in cold regions has an important impact on the durability of concrete structures. Freeze–thaw damage to concrete in alpine regions of China is a common issue, and it poses a huge threat to the safety of building use and economic and environmental protection. Frost resistance is particularly important. The main reason for the freeze–thaw damage of concrete is that the water in the concrete pores, under the action of alternating dry and wet and freeze–thaw cycles, forms fatigue stress from the combined action of ice expansion pressure and osmotic pressure, which causes the concrete to produce denudation damage from the surface to the inside, thereby reducing the strength of the concrete [5–7]. Research on the antifreeze performance of concrete structures under low-temperature freeze–thaw cycle environment can not only reveal potential dangers and avoid major safety accidents but also provide a basis for

**Citation:** Tan, Y.; Long, J.; Xiong, W.; Chen, X.; Zhao, B. Effects of Polypropylene Fibers on the Frost Resistance of Natural Sand Concrete and Machine-Made Sand Concrete. *Polymers* **2022**, *14*, 4054. https:// doi.org/10.3390/polym14194054

Academic Editors: R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 18 August 2022 Accepted: 15 September 2022 Published: 27 September 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

durability design, testing, and reinforcement of concrete structures in low-temperature freeze–thaw environments. Improving the frost resistance of concrete has become a hot topic of research at home and abroad [8–10].

Compared with ordinary concrete, adding fiber is one of the methods that can effectively reduce the brittleness and improve the toughness of concrete [11–14]. Fiber materials have been widely studied by scholars at home and abroad. The results show that fiber structures have the characteristics of strong plasticity, high toughness, and strong adhesion to concrete, and adding fibers into concrete can offset the internal stress and play an effective role in frost resistance [15–18]. The different types of fibers have different effects on the frost resistance of concrete. Among them, not only is polypropylene fiber (PPF) small in diameter, low in quality, low in cost, and good at self-dispersion, but it also has the characteristics of inhibiting concrete plastic cracking, preventing crack propagation, limiting matrix damage, improving the long-term working performance of concrete structures, and improving structural durability [19–21]. In the process of freezing, the increase in the elastic modulus of PPF can effectively offset the ice expansion force, and the decrease in the elastic modulus of PPF during the thawing process helps it to release the accumulated expansion energy. Therefore, the change in the elastic modulus of PPF is enhanced as a whole. The strain capacity of concrete in the freeze–thaw environment can effectively inhibit the frost heave cracking of concrete. Different lengths and dosages of PPF have varying degrees of influence on the tensile properties, fatigue resistance, and wear resistance of concrete [22–24]. The performance of concrete has been improved in all aspects [25]. Cheng Hongqiang et al. [26] conducted a freeze–thaw damage test of polypropylene fiber concrete. Under the action of freeze–thaw cycles, the damage of polypropylene fiber concrete continued to accumulate, and the mass loss rate and split tensile strength continued to decrease. The splitting tensile strength has been continuously improved. Chen Liuzhuo et al. [27] studied the frost resistance of ordinary concrete by adding steel fibers and polypropylene fibers. The results show that controlling the amount and length of fibers can effectively improve tensile strength and flexural strength. Durability is enhanced, and polypropylene fibers are more damage-resistant than steel fibers. Salehi Parisa et al. [28] studied the effect of PPF on the mechanical properties and durability of reinforced lightweight concrete through the mixed design of three volume dosages (0%, 0.5%, and 1%) of PPF, and the results show that the incorporation of fibers can improve compressive, tensile, and flexural strength, and it also reduces the water absorption and permeability of concrete. Liu Bo et al. [29] analyzed and summarized the influence of polypropylene fibers on the mechanical properties and durability of concrete and proved that the addition of polypropylene fibers can make concrete mechanically resistant to compression, tension, bending, and impact. In terms of performance and durability, including frost resistance, impermeability, and carbonation resistance, it is superior to ordinary concrete.

Due to the over-exploitation of natural sand, natural sand resources are depleted, which seriously affects the ecological environment, and finding new alternative sources has become a hot topic. Machine-made sand is artificially mined industrial waste slag, tailings, etc., which are mechanically crushed and screened into rocks with a particle size of less than 4.75 mm [30–32]. The mud powder in natural sand has an adverse effect on the working performance, volume stability, and durability of concrete, while the finegrained stone powder in machine-made sand can improve the gradation of machine-made sand and fill the pores of the particles. A large number of studies have shown that an appropriate amount of stone powder can improve the workability of concrete and improve the strength and durability of machine-made sand concrete [33–35]. Replacing natural sand with manufactured sand has different degrees of influence on the mechanical properties of concrete. Ding et al. [36] prepared manufactured sand concrete with a compressive strength of higher than 50 MPa, and some scholars have prepared ultra-high-pressure concrete through different treatment methods. The compressive strength of high-performance concrete exceeds 140 MPa [37–40]. Bonavetti and Irassar [41] studied the effect of machinemade sand and gravel powder on the properties of mortar. With an increase in age and in

the replacement rate of machine-made sand, the compressive strength of concrete showed an increasing trend. Donza et al. [42] studied the effects of several different materials of manufactured sand and natural sand on the properties of high-strength concrete. The results showed that the mechanical properties of high-strength concrete were better, due to the mechanical occlusion of the manufactured sand. Zeghichil et al. [43] studied the effect of natural sand and multi-angled machine-made sand on the working performance and mechanical properties of self-compacting concrete. Multi-angled machine-made sand requires more water to achieve the same working performance, but the compressive strength of the concrete is lower than that of regular concrete. The flexural strength is higher than that of concrete prepared from natural sand.

Regarding the application of fiber materials in machine-made sand concrete, Kandasamy and Murugesan [44] studied the compressive strength and split tensile strength of machine-made sand concrete and concrete mixed with plastic fibers. The results show that the addition of plastic fibers can effectively improve the mechanical properties and durability of manufactured sand concrete. Some scholars have improved the shortcomings of easy cracking and poor frost resistance by admixing PPF and mineral admixtures in desert sand concrete [45]. In view of the excellent performance of polypropylene fibers in concrete compared with other fibers, there are relatively few studies on the freeze–thaw cycle of polypropylene fiber concrete in which continuous graded machine-made sand replaces 100% of natural fine aggregates.

In this paper, by incorporating polypropylene fibers of different lengths and volumes, natural sand concrete and manufactured sand concrete with a replacement rate of 100% were used as the objects to study the change law of the durability of PFNSC and PFMSC under freeze–thaw cycles. A fitting model corresponding to the fiber lifting rate and contribution rate [46] was established to quantitatively characterize the influence of PPF. At the same time, based on learning from the classical combination model of damage [47–51], the damage theory under freeze–thaw cycles was used [52,53]. A freeze–thaw damage model and a response surface model were established to analyze the frost resistance of PFNSC and PFMSC, and the effects of fiber content, fiber length, and freeze–thaw times on the frost resistances of PFNSC and PFMSC were studied. The PFNSC and PFMSC specimens were compared.

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

#### *2.1. Raw Materials and Mix Ratio*

The cementitious materials used in this test are cement and grade I fly ash, the cement type is standard P.O. 42.5 grade ordinary cement, and all the cement indexes meet the requirements of GB175-2007 "General Portland Cement". Among the materials, fly ash complies with the requirements of the standard GB/T 1596-2005 "fly ash used in cement-based concrete", its apparent density is 2500 kg/m<sup>3</sup> , the specific surface area is 455.2 m2/kg, and it is a 30 mm continuous graded natural aggregate; fine aggregates are natural river sand (NRS) and machine-made sand(MS) in zone II, for which the fineness modulus is 2.9, the apparent density of natural river sand is 2650 kg/m<sup>3</sup> , and the apparent density of manufactured sand is 2610 kg/m<sup>3</sup> . According to the standard GB/T14684-2011 "Sand for Construction", the MB value and stone powder content of machine-made sand meet the requirements of Class II machine-made sand. The superplasticizer is polycarboxylate, and the gradation curve is shown in Figure 1. The physical map and microscopic topography of the polypropylene fibers are shown in Figure 2a,b, and their performance indicators are shown in Table 1.

**Figure 1.** Cumulative sieve residue of fine sand.

**Figure 2.** Physical and microscopic pictures of polypropylene fibers: (**a**) physical map, (**b**) micrograph.



It can be seen from Figure 1 that the surface characteristics of the gradations of machinemade sand are that the occupancy of the middle section is higher than the occupancy of both ends, that is, "large in the middle and small at both ends", and the degree of gradation continuity is poorer than that of natural fine aggregate. In this test, the benchmark concrete strength grade is C30, and natural river sand and machine-made sand are used as fine aggregates. The mix ratio of fiber concrete is shown in Table 2. The percentage of polypropylene fibers should be controlled at about 1.0% and not more than 1.2% [54]. In this paper, the volume content of PPF is 0%, 1.0%, and 1.2%. PFNSC a–b and PFMSC a–b are 12 mm and 19 mm ("a" is the volume content of PPF and "b" is the length of PPF), respectively.


**Table 2.** Laboratory mix proportion of concrete.

#### *2.2. Test Equipment and Test Methods*

A total of 14 sets of specimens were designed for the test, all of which conform to GB/T 50082-2009 "Standards for Test Methods for Long-term Performance and Durability of Ordinary Concrete".

The compressive test uses a DYE-2000A microcomputer servo pressure-testing machine (produced by Cangzhou Zhulong Engineering Instrument Co. Ltd., Cangzhou, China), as shown in Figure 3a, and a cube with a size of 100 mm × 100 mm × 100 mm to measure the compressive strength test. The test loading speed is 0.5 MPa/s, and the compressive strength correction coefficient is 0.95. After the test piece reaches the test age, it is taken out of the curing room. The surface of the test piece and the upper and lower bearing plates are wiped, and a check is made for whether the center of the test piece is aligned with the center of the lower pressure plate of the testing machine. After the inspection is completed and confirmed to be correct, the compressive strength test is carried out on the testing machine. The loading speed of the press is manually controlled. When the specimen is close to failure and begins to deform sharply, we stop adjusting the test throttle until failure and record the failure load. The arithmetic mean of the measured values of 3 test pieces is taken as the strength value of the test piece.

The split tensile test uses a DYE-2000A microcomputer servo pressure-testing machine (produced by Cangzhou Zhulong Engineering Instrument Co., Ltd., Cangzhou, China), as shown in Figure 3b. Cubes with dimensions of 100 mm × 100 mm × 100 mm were tested for split tensile strength. The test loading speed is 0.08 MPa/s. After the specimen is taken out, the dry and wet state of the specimen are kept unchanged, and measurements are taken to determine whether the size of the specimen meets the requirements. A line is drawn in the middle of the side of the specimen during forming to determine the position of the splitting surface, the lower cushion and gasket are placed in the center of the lower plate, and the specimen, upper gasket, and cushion are placed in sequence on the lower gasket. The contact busbar of the gasket is aligned with the load contact line on the specimen accurately, and the contact between the upper pressure plate and the cushion is adjusted. After checking the adjustment and correctness, turn on the split tensile testing machine. When the specimen is close to failure or deformation, we stop adjusting the test throttle until failure, and the value is recorded. Three test blocks are set for each group of mix ratios, and the value method is the same as that of the compressive strength value.The freeze–thaw

test is performed using a TDR-III(produced by Hebei Huawei Co., Ltd., Cangzhou, China), automatic rapid freeze–thaw test machine, as shown in Figure 3c, with a cylindrical block with a size of 100 mm × 100 mm × 400 mm, and a freeze–thaw cycle test. This is carried out according to the "slow freezing method" in GB/T 50082-2009 "Standard for Long-term Performance and Durability Test of Ordinary Concrete". After the test piece reaches the test age, it is taken out from the curing room four days in advance and immersed in water with a temperature of 15~20 ◦C. After soaking, the test piece is taken out to dry the surface moisture, and its initial mass is weighed. A total of 150 freeze–thaw cycles were carried out in this test, and measurements after every 50 freeze–thaw cycles were used to collect relevant frost resistance durability parameters. The collected contents included: concrete specimen quality, compressive strength, and flexural strength.

**Figure 3.** Testing device: (**a**) Compression test, (**b**) Flexural strength test, (**c**) Freeze–thaw cycle test.

#### **3. Test Results and Discussion**

#### *3.1. Strength Analysis after Freeze–Thaw Cycle*

The PFNSC and PFMSC specimens were set up in the experimental group and the control group with different fiber content, fiber length, and manufactured sand content. They were divided into 14 groups, and each group was set with three specimens, a total of 42 test blocks. The compressive strength, split tensile strength, and quality changes after 0, 50, 100, and 150 freeze–thaw cycles are shown in Table 3.

**Table 3.** Compressive strength, split tensile strength, and quality under different freeze–thaw cycles.


Using the results of the orthogonal test, the compressive strengths of the PFNSC and PFMSC specimens with different freezing and thawing times and fiber lengths under different polypropylene contents were generated using origin 2021 software (produced by OriginLab Co. Ltd., Northampton, MA, USA), as shown in Figure 4. It can be seen from Figure 4a,b that, with the increase in the number of freeze–thaw cycles, different PPF volume contents in PFMSC specimens had a greater impact on the compressive strengths of the specimens. With an increase in the number of freeze–thaw cycles, the strength of the specimen with a volume of 1.0% decreased rapidly. The strength of the concrete specimen with a volume of PPF of 1.2% was higher than that of the specimen with a 1.0% volume of PPF under different freeze–thaw cycles. The 1.0% concrete specimen shows that the 1.2% fiber volume content results in better frost resistance than the 1.0% fiber volume content. After 50–150 freeze–thaw cycles, the compressive strength of the specimen decreased first, then increased, and then decreased with the increase in fiber length, and the deceleration rate reached the maximum when the fiber length was 12 mm. The comparison of Figure 4a–d shows that with the increase in freeze–thaw times, the average reduction rate for the compressive strengths of the PFNSC and PFMSC specimens is close. The average reduction in the compressive strengths of PFNSC specimens is 26.4%; the average reduction in the compressive strengths of the PFMSC specimens is 26.2%. This shows that the frost resistance of the PFNSC and PFMSC specimens is equivalent.

Figure 5 shows the split tensile strength of PFNSC and PFMSC specimens with different freezing and thawing times and fiber lengths, respectively, for different polypropylene contents. It can be seen from Figure 5a,b that after 50–150 freeze–thaw cycles of the PFMSC specimens, the split tensile strength of the concrete is significantly lower than that of the control group without PPF. The split tensile strength decreases with the increase in fiber length, showing a trend of increasing first and then decreasing. For example, the split tensile strength of the PFMSC specimen with 50 freeze–thaw cycles and 1% PPF volume content is 1.47 times that of the plain concrete without a PPF resemblance. When the volume content of PPF is constant, the relative split tensile strengths of the concrete specimens gradually decrease with the increase in the number of freeze–thaw cycles, but after 100 freeze–thaw cycles, the split tensile strengths of PFMSC specimens decreases. Under the condition of constant fiber content, the relative split tensile strength decreased with the increase in fiber

length, with the strength first increasing and then decreasing. By comparing Figure 5a–d, it is found that with the increase in freezing and thawing times, the average reduction rate of the split tensile strengths of PFNSC specimens is slightly higher than that of PFMSC specimens, indicating that the replacement of natural sand concrete by machine-made sand concrete can effectively enhance the durability of concrete. This is because appropriate stone powder in the PFMSC specimen can improve the grading of machine-made sand and fill the pores of particles, which improves the durability of concrete. After 50–150 freeze– thaw cycles, the average reduction in the splitting tensile strengths of PFNSC specimens was 32.5%; the average reduction in splitting tensile strengths of the PFMSC specimens was 23.2%. This shows that the antifreeze performance of PFMSC is better than that of the PFNSC specimens.

**Figure 4.** Compressive strength test value at different freeze–thaw cycles: (**a**) PFMSC compressive strength value when PPF volume is 1.0%, (**b**) PFMSC compressive strength value when PPF volume is 1.2%, (**c**) PFNSC compressive strength value when PPF volume is 1.0%, (**d**) PFNSC compressive strength value when PPF volume is 1.2%.

**Figure 5.** Split tensile strength at different freeze–thaw cycles: (**a**) PFMSC splitting tensile strength value when PPF volume is 1.0%, (**b**) PFMSC splitting tensile strength value when PPF volume is 1.2%, (**c**) PFNSC splitting tensile strength value when PPF volume is 1.0%, (**d**) PFNSC splitting tensile strength value when PPF volume is 1.2%.

The test results show that freeze–thaw cycles have a great influence on the compressive strengths and split tensile strengths of the PFNSC and PFMSC specimens. The addition improved the antifreeze properties of PFNSC and PFMSC. Different PPF volume contents and PPF lengths also have a certain degree of influence on the mechanical properties of concrete. It can be seen that, after 0–150 freeze–thaw cycles, when the volume of PPF is the same, the durability energies of PFMSC and PFNSC increase first and then decrease with the increase in the length of PPF; when the length of PPF is 6–12 mm, the mechanical properties of PFMSC and PFNSC increase with the increase in fiber volume content. When the PPF length is 19 mm, the mechanical properties of PFMSC and PFNSC decrease with the increase in fiber volume content.

Figure 6 shows the apparent state of the PFMSC specimen after 0, 50, 100, and 150 freeze–thaw cycles and the surface macro-state of the binarized concrete specimen under the same variables. It can be seen from Figure 6 that after 0 freeze–thaw cycles, there is a small area of cement material shedding on the surface of the specimen; that is, the black area marked in red is the area where a small part of the surface of the specimen has

fallen off the cement material. With the increase in the number of freeze–thaw cycles, the quality of the PFMSC specimen continued to decrease, and the coarse aggregate inside the PFMSC specimen was gradually exposed. After 50 freeze–thaw cycles, a small area of the surface of the specimen had fallen off, and the mass loss was approximately 0.32%. After 100 freeze–thaw cycles, much of the surface of the test piece had fallen off, large coarse aggregate particles were exposed, and the mass loss was approximately 1.85%. When the freeze–thaw cycles reached 150, the surface of the specimen showed a covered black area, indicating that the surface of the concrete specimen was almost completely peeled off under 150 freeze–thaw cycles, the aggregate under the surface was completely exposed, and the mass loss was approximately 4.5%. With the increase in the number of freeze–thaw cycles, the surface peeling degree of the specimens gradually deepened, the strength damage of the PFMSC specimens was serious, and the mass loss was large.

**Figure 6.** Binarization processing of specimen images with different freezing and thawing times. (**a**) Concrete apparent structure before and after freeze and thawing. (**b**) Binary treatment diagram before and after freeze and thawing.

#### *3.2. Freeze–Thaw Injury of PFNSC and PFMSC under Strength Evaluation Index*

In order to consider the change in freeze–thaw cycles and the decay law of concrete mechanical properties, the relationship between the compressive and split tensile strengths of the freeze–thaw cycle damage degree under different polypropylene fiber volume contents and lengths was analyzed. Based on the data changes in compressive strength and splitting tensile strength, according to "Concrete Damage Mechanics", the degree of freeze–thaw cycle damage of concrete is represented by *D*, and *D*<sup>c</sup> and *D*<sup>t</sup> are defined as the degree of freeze–thaw cycle damage under compression and splitting resistance, respectively, which correspond to strength index damage. They are calculated by Formulas (1) and (2).

$$D\_{\rm c,m} = 1 - \frac{f\_{\rm c,n}}{f\_{\rm c,0}} \tag{1}$$

$$D\_{\rm t,m} = 1 - \frac{f\_{\rm t,n}}{f\_{\rm t,0}} \tag{2}$$

In the formulas:

*D*c,m—Compressive strength damage variable (%); *D*t,m—Split tensile strength damage variable (%);

*f*c,n—Compressive strength (MPa) under corresponding freeze–thaw times; *f*t,n—Splitting strength (MPa) under corresponding freeze–thaw times.

In the formula, 'c' indicates that the specimen is a compression specimen, 't' indicates that the specimen is a split-pull specimen, 'n' is the number of freeze–thaw cycles, and 'm' is the content of manufactured sand. Through Formulas (1) and (2), the strength damage model of the polynomial function under the freeze–thaw cycles was established, and the

strength damage of the PFNSC and PFMSC specimens was predicted using the model, as shown in Figure 7a–d under different freeze–thaw cycles.

**Figure 7.** Strength damage under different freeze–thaw cycles: (**a**) PFNSC compressive strength damage rate, (**b**) PFMSC compressive strength damage rate, (**c**) PFNSC splitting tensile strength damage rate, (**d**) PFMSC splitting tensile strength damage rate.

It can be seen from Figure 7 that the extents of freeze–thaw damage of the compressive strengths and split tensile strengths of the PFNSC and PFMSC specimens increase with the increase in the number of freeze–thaw cycles. From Figure 7a,b on tensile strength damage, it can be seen that compared with PFNSC, the compressive strength of the PFMSC specimen is more damaged, because the shape of the machine-made sand is sharp and rough and has many edges and corners. After a certain number of freezing and thawing cycles, the machine-made sand will accelerate the material falling off. After PFNSC freeze–thaw cycles, the freeze–thaw damage rate without PPF is greater than that with PPF. This is because the polypropylene fibers are constrained in the concrete and bear part of the damage caused by the freeze–thaw cycles inside the specimen. With the effect of force, the freeze–thaw damage rate of PFMSC without PPF is lower than the freeze–thaw damage rate of the PPF-doped specimens, and generally, the longer the length of PPF, the greater the amount of freeze–thaw damage. This is due to the mechanism of the sand particles. The influence of shape characteristics on the compressive strengths of PFMSC specimens is greater than

that of fibers, and the longer the fibers, the larger the contact area with machine-made sand will be. It can be seen from Figure 7c,d that the split tensile strength damage rate is significantly reduced by PPF. The freeze–thaw damage rate of the specimens without PPF is greater than the freeze–thaw damage rate of the specimens with PPF. Longitudinal pressure inside the specimen can improve the durability of concrete; PFNSC specimens have the same length of PPF, and the splitting tensile strength of the specimen when the volume content of PPF is 1.2% is less than that of the specimen when the volume content of PPF is 1%. The freeze–thaw damage rate of PFMSC is due to the bending of PPF when the inside of the specimen is greatly constrained, which increases the contact area and improves the anti-splitting ability; PFMSC specimens have the same length of PPF, and the volume content of PPF is 1.2%. The freeze–thaw damage rate of the specimen is greater than the freeze–thaw damage rate of the specimen with a volume content of 1% PPF. The mechanism is that the influence of the shape characteristics of the machine-made sand particles on the splitting tensile strength of PFMSC is greater than that of the fiber inside the specimen. The fitting function of the freeze–thaw damage rate is shown in Tables 4 and 5. It can be seen from Tables 4 and 5 that the fitted regression curve and the observed value of R <sup>2</sup> are both greater than 0.95, indicating that the fitted model is highly reliable.

**Table 4.** Fit function of compressive strength damage value.


**Table 5.** Fit function of split-pull strength damage values.


## **4. Model of the Composite Factor RSM Intensity**

The RSM uses mathematical and statistical methods to model and analyze problems affected by multiple variables, with the ultimate goal of optimizing the response value [55,56]. Box and Wilson first proposed the response surface method. At that time, the research on the response surface method was limited to how to obtain an explicit function using statistical methods to approximate a complex implicit function. Fang et al. [57] used the

D-optimal design and a first-order response surface model to predict the dynamic response and damage identification of intact and damaged systems, and they also used numerical examples, reinforced concrete frame model tests, and I-40 real bridge test results to verify the effectiveness of the proposed method. Using the response surface method (RSM) to model the relationship between factors and levels can better analyze the accuracy, significance, and reliability of the experimental data. Zong Zhouhong et al. [58] used the center composite design (CCD) method and the response surface model to complete the finite element model correction of the Baishi Bridge, and they proved that the bridge finite element model correction based on the response surface method has a higher accuracy. Many scholars at home and abroad have established relational models through the response surface method and have obtained correspondingly optimal results [59–62].

The damage degree of PFNSC and PFMSC under different conditions was quantitatively analyzed via freeze–thaw damage. This showed that when the length of PPF is the same and the volume content of PPF is 1.2, the antifreeze performance of PFNSC is optimal; when the volume content of PPF is 1.0, the antifreeze performance of PFMSC is optimal. In order to better study the effect of PPF length and the number of freeze–thaw cycles on antifreeze performance, that is, taking the PPF length and the number of freeze–thaw cycles as two factors, a response surface strength model was established, and design-Expert software was used to conduct a multivariate analysis of the experimental data. Regression analysis was performed to establish the fitting model of Equations (3)–(6). Equations (3)–(6) were analyzed via variance analysis. The regression coefficient value *R* <sup>2</sup> was used to test the reliability of the model. When the regression coefficient value *R* <sup>2</sup> was close to 1, the model reliability was high.

$$f\_{\text{f0}} = 63.91 - 2.63A + 0.13B + 2.95C + 1.59AB - 0.81AC - 2.36BC + 2.13A^2 - 8.27B^2 + 5.4C^2 \tag{3}$$

$$R^2 = 0.8851$$

$$f\_{\rm c100} = 46.57 - 2.56A - 0.46B - 0.75C - 0.97AB - 1.24BC + 3.25A^2 - 2.8B^2 - 7.04C^2 \tag{4}$$

$$R^2 = 0.9050$$

$$f\_{l0} = 2.4 - 0.47A - 0.083B + 0.37C + 0.027AB - 0.1BC + 0.085A^2 - 0.17B^2 + 0.22C^2 \tag{5}$$

$$R^2 = 0.9368$$

$$f\_{100} = 3.57 - 0.36A + 0.011B + 0.28C + 0.12AB + (2.85E - 003)BC - 0.27A^2 - 0.018B^2 - 0.59C^2 \tag{6}$$

$$R^2 = 0.9248$$

In the formula:

*f*c0—PFNSC compressive strength (MPa); *f*c100—PFMSC compressive strength (MPa); *f*t0 —PFNSC tensile strength (MPa); *f*t100—PFMSC split tensile strength (MPa);

*A*—Freeze–thaw times (*n*); *B*—Polypropylene fiber length (mm); *C*—Polypropylene fiber volume content (%).

It can be seen from the fitting model that the values obtained by the fitting equation are all close to 1 and that the value of the coefficient of variation obtained by the model is very small, indicating that the correlation between the model and the test data is significant and the degree of fitting is high, which indicates the model can better analyze and predict freezing conditions. The effects of the PPF length and the number of freeze–thaw cycles on the antifreeze properties of PFNSC and PFMSC specimens within 150 thaw cycles are found. Figures 8 and 9 show the optimal three-dimensional response surfaces and contour maps of the compressive and split tensile strengths of the PFNSC specimen and the PFMSC specimen obtained, respectively, based on the RSM model within 150 freeze–thaw cycles.

**Figure 8.** Effects of Factor A, B, and their interactions on R.(PFNSC): (**a**) Surface diagrams of compressive strength variation with A and B (2D), (**b**) Surface diagrams of compressive strength variation with A and B (3D), (**c**) Surface diagrams of split tensile strength variation with A and B (2D), (**d**) Surface diagrams of split tensile strength variation with A and B (3D).

When the volume content of PPF is 1.2%, the PFNSC contour map and 3D response surface are as shown in Figure 8. Through the analysis of Figure 8a,b, it can be seen from the contour map that the green on both sides gradually changes to orange in the center and that the center is the region with the highest responsivity, indicating that the material content corresponding to this region is the optimal doping value. In the three-dimensional response surface graph, the variation trend of the R value along factor B is greater than that of factor A, indicating that factor B (number of freeze–thaw cycles) has a greater impact on the R value than factor A (fiber length). When the number of freeze–thaw cycles increased from 50 to 150, the compressive strength and split tensile strength of PFNSC decreased by 27.5 MPa and 1.8 MPa, respectively; that is, the strength of the specimen decreased with the increase in the number of freeze–thaw cycles. When the fiber increased from 6 mm to 19 mm (1.8 MPa), the strength of the specimen showed a trend of first increase and then decrease, and when the fiber length was approximately 10 mm, the tensile strength reached the maximum. As can be seen from Figure 8c,d, the green on the top of the contour map gradually changes to the red on the bottom. The region below the center is the region with the highest response, and the corresponding polypropylene fiber length and the number of freeze–thaw cycles reach optimum values. The density of the contour lines on the ordinate in the split tensile strength contour diagram is greater than the density of the contour lines on the abscissa; that is, when the fiber volume content is 1.2%, the effect of fiber length on

the split tensile strength is greater than that of freezing. The effect of the number of melting cycles is seen.

**Figure 9.** Effects of Factor A, B, and their interactions on R.(PFMSC): (**a**) Surface diagrams of compressive strength variation with A and B (2D), (**b**) Surface diagrams of compressive strength variation with A and B (3D), (**c**) Surface diagrams of split tensile strength variation with A and B (2D), (**d**) Surface diagrams of split tensile strength variation with A and B (3D).

As shown in Figure 10, when the volume content of the PFNSC fibers is 1.2%, the response surface model can predict that when the number of freeze–thaw cycles is 105 and the fiber length is 11.8 mm, the compressive strength and splitting tensile strength of PFNSC are both at maximum values of 33.8 MPa and 3.1 MPa, respectively, indicating that PFNSC can maintain a good antifreeze performance within 105 freeze–thaw cycles. The desirability is 0.975, indicating that the model has high prediction reliability.

As shown in Figure 11, when the volume content of the PFMSC fibers is 1.0%, the response surface model predicts that when the number of freeze–thaw times is 96 and the fiber length is 9.1 mm, the compressive strength and splitting tensile strength of PFMSC are both at maximum values of 41.21 MPa and 3.2 MPa, respectively, indicating that PFMSC can maintain a good antifreeze performance within 96 freeze–thaw cycles. The desirability is 0.826, indicating that the model has high prediction reliability.

**Figure 10.** PFNSC and optimization results of strength maximization.

**Figure 11.** PFMSC and Optimization results of strength maximization 20.

## **5. Micromorphology Analysis**

Scanning electron microscopy (SEM) using a Czech TESCAN MIRA LMS was used to observe the microscopic morphologies of manufactured sand concrete specimens without PPF and PFMSC specimens, before and after the freeze–thaw cycle test. The microscopic morphologies of the PFMSC0–0 specimen and the PFMSC specimen before and after freezing and thawing are shown in Figure 12. After freezing and thawing of PFMSC0–0, the matrix pores increase and increase, and deep cracks appear; the surface of the PFMSC specimen before the freezing and thawing test is relatively smooth and flat, with fewer cracks and pores. More cracks and pores appear on the surface of the matrix after the freeze– thaw test, but because PPF belongs to the class of bundled monofilament organic fibers, it is distributed in three-dimensional random directions in concrete, and a network-like reinforcement system is formed inside it. The cracking of concrete cracks and the expansion of cracks caused by drying and chemical shrinkage during the cement hydration process are inhibited. After the freeze–thaw cycle, the connection between the fibers and the matrix is still relatively tight, indicating that the addition of PPF can effectively improve the frost resistance of manufactured sand concrete.

**Figure 12.** PFMSC microtopography of PFMSC before and after freeze–thawing: (**a**) PFMSC0–0: Before freezing and thawing, (**b**) PFMSC0–0: After freezing and thawing, (**c**) PFMSC1.0–9: Before freezing and thawing, (**d**) PFMSC1.0–9: After freezing and thawing.

## **6. Conclusions**


damage of the PFMSC specimen is lower; when the volume content of PPF is 1.0% and the length is 6 mm, the splitting tensile strength damage of the PFMSC specimen is higher. The strength damage of the PFMSC specimens is generally lower than that of the PFNSC specimens, and when the PPF length is the same, the volume content of 1% can better reduce the strength damage. According to the prediction results of PFNSC and PFMSC, the antifreeze performances of PFNSC and PFMSC are similar. This shows that polypropylene fibers have similar effects on PFNSC and PFMSC. It also shows that it is feasible to replace natural sand concrete with 100% artificial sand in practical engineering.

4. The optimal performances of PFNSC and PFMSC are predicted by the RSM strength composite model. PFNSC can maintain a good antifreeze performance within 105 cycles of freezing and thawing. When the volume content of PPF is 1.2% and the length is 11.82 mm, the freezing performance is optimal, the compressive strength value is 33.8 MPa, and the split tensile strength value is 3.1 MPa. PFMSC can maintain a good antifreeze performance within 96 freeze–thaw cycles. When the volume content of PPF is 1.2% and the length is 9.1 mm, the antifreeze performance of the specimen reaches its maximum, its maximum tensile strength value is 45.8 MPa, and the split tensile strength value is 3.2 MPa.

**Author Contributions:** Conceptualization, Y.T., J.L., X.C., W.X. and B.Z.; methodology and data curation, Y.T. and J.L.; validation and investigation, Y.T., J.L. and B.Z.; formal analysis, Y.T., J.L.; resources, B.Z.; writing—original draft preparation, Y.T.; writing—review and editing, J.L. and B.Z.; supervision, X.C.; and funding acquisition, X.C. All authors have read and agreed to the published version of the manuscript.

**Funding:** The central government guides local science and technology development projects (No. 2020CFA046); Demonstration Construction Project of Ecological Science and Technology Manor Base (No. 2018ZYYD037).

**Data Availability Statement:** The data used to support the findings of this study are available from the corresponding author upon request.

**Acknowledgments:** The authors acknowledge the Hubei University of Technology for funding this research.

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

## **References**


## *Article* **Reinforced Structure Effect on Thermo-Oxidative Stability of Polymer-Matrix Composites: 2-D Plain Woven Composites and 2.5-D Angle-Interlock Woven Composites**

**Xingzhong Gao 1,2, Tiancong Han 1,2 , Bolin Tang <sup>3</sup> , Jie Yi <sup>3</sup> and Miao Cao 3,\***


**Abstract:** The thermo-oxidative stability of carbon fiber polymer matrix composites with different integral reinforced structures was investigated experimentally and numerically. Specimens of 2-D plain woven composites and 2.5-D angle-interlock woven composites were isothermally aged at 180 ◦C in hot air for various durations up to 32 days. The thermal oxidative ageing led to the degradation of the matrix and the fiber/matrix interface. The degradation mechanisms of the matrix were examined by ATR-FTIR and thermal analysis. The interface cracks caused by thermal oxidative ageing were sensitive to the reinforced structure. The thermo-oxidative stability of the two composites was numerically compared in terms of matrix shrinking and crack evolution and then experimentally validated by interlaminar shear tests.

**Keywords:** thermo-oxidative stability; woven composites; structure effect; finite element

## **1. Introduction**

Advanced textile composites are gaining market share in various industries, including examples in the aerospace [1], maritime [2], automotive [3], civil infrastructure [4,5], and wearable electronics [6] industries, due to their exceptional electrical, mechanical, and thermal properties. Woven composites are becoming a research hotspot as they are one of the most advanced textile composites [7–9] and have shown great potential for aircraft applications, such as wings and engine blades, due to the high strength/weight ratio and impact resistance [10–12]. Aerospace applications require a long service life of materials while the thermal oxygen ambient will always be met during the service, which causes the reduction of composite properties and threatens aircraft safety [13].

The polymer matrix is susceptible to temperatures. The thermal properties of the polymer matrix have been widely reported by previous work [14,15] using dynamic mechanical analysis, thermogravimetric analysis, differential scanning calorimeter analysis, and so on. The obtained glass transition temperature or decomposition temperature could be regarded as an important indicator for thermal stability. The elevated temperature can also accelerate the oxidation rate, and the combined effect of thermolysis and oxidation promotes chain scission, accompanied by the departure of low-molecule volatiles, and finally leads to chemical shrinkage [16–18]. The matrix shrinkage has been quantitatively measured at a microscopic scale by many scholars [19–22]. While in the fiber reinforced polymer composites, most of the reinforcement, such as carbon fiber, is reasonably inert to thermal oxidation in service (<250 ◦C) [23]. Thus, the matrix shrinkage causes a mismatch of deformation between the fiber and matrix and then induces a tensile stress within the

**Citation:** Gao, X.; Han, T.; Tang, B.; Yi, J.; Cao, M. Reinforced Structure Effect on Thermo-Oxidative Stability of Polymer-Matrix Composites: 2-D Plain Woven Composites and 2.5-D Angle-Interlock Woven Composites. *Polymers* **2022**, *14*, 3454. https:// doi.org/10.3390/polym14173454

Academic Editors: Rushdan Ahmad Ilyas, Salit Mohd Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 21 July 2022 Accepted: 18 August 2022 Published: 24 August 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

fiber–matrix interfacial phase [24]. The stress accumulates with ageing time and eventually leads to interface cracks, which create additional pathways for oxygen diffusion and compromise the structural integrity of composite materials [25,26].

The distribution of interface cracks shows the relationship of the architectural structure of fibers. For composite laminates, the ply orientation angles [27], fiber spacing [28,29] as well as stacking sequences [30] have a significant influence on the initiation and propagation of the cracks. The evolution of ageing cracks is much harder to predict in fabric reinforced composites due to the complex interlacing structures. Anisotropic distribution of ageing cracks has been reported in braided and woven composites due to different yarn spatial configurations [31–34]. The thermal oxidative stability of the composites could be affected by the reinforced structure accordingly. Wu [35] and Fan [36] compared the ageing properties of braided composites with unidirectional/laminated composites, respectively, and both results indicated that the braided composites could improve durability thanks to better structural integrity. Revealing the effect of reinforced structure on the thermo-oxidative stability of composites has good benefits for the durability design of the composite structure.

Most of the previous research has mainly focused on the ageing behaviors of composites with some specific reinforced structures. Only a few papers investigated the reinforced structure effect by comparing the degradation of overall mechanical properties. How the reinforced structure affects the distribution of matrix shrinkage and interface cracks has not yet been revealed. Herein, we present a comprehensive investigation into the thermaloxidative degradation mechanism of 2-D plain woven composites (2-D PWC) and 2.5-D angle-interlock woven composites (2.5-D AWC). The thermal degradation mechanisms of the matrix were evaluated by chemical and thermal analysis. The reinforced structure effect on the thermal-oxidative stability of the two woven composites was explored by both numerical and experimental approaches. The structure effect was first compared by the distribution of shrinkage displacement as well as the evolution of interface cracks, and then validated by interlaminar shear tests.

#### **2. Experimental**

#### *2.1. Material*

In this paper, a diglycidyl ether of bisphenol-A (DGEBA) epoxy resin (JC-02A/JC-02B, Changshu Jaffa Chemical Inc., Suzhou, China) was selected as the resin matrix. The carbon fiber tows were provided by Toray Industries, Inc., Japan. The specifications of the epoxy matrix and carbon fiber were listed in Tables 1 and 2, respectively.

**Name Chemical Component Viscosity at Room Temperature (MPa**·**s) Epoxide Number (eq/100 g) Blending Ratio** JC-02A Bisphenol A epoxy resin 1000–3000 0.5–0.53 100:80 JC-02B Modified anhydride 30–50 -

**Table 1.** Specifications of the epoxy matrix.

**Table 2.** Specifications of the carbon fiber.


#### *2.2. Material Preparation*

Figure 1 shows the structures of 2-D plain woven fabric and 2.5-D angle-interlock woven fabric. The specifications of the preforms are listed in Table 3. Epoxy resin was used

to impregnate the fabric with the vacuum assisted resin transfer method (VARTM). The composite was consolidated with the curing process: 90 ◦C for 2 h, 110 ◦C for 1 h, and 130 ◦C for 4 h in sequence, and the vacuum was about 0.1 MPa. As the single plain woven fabric ply is very thin, the 2-D plain woven composite was obtained by stacking 20 plies of fabrics along the thickness direction (in a 0◦ direction). The fiber volume fraction is 44.6% and 45.2% for 2.5-D angle-interlock woven composite and 2-D plain woven composite, respectively, obtained by muffle furnace combustion. used to impregnate the fabric with the vacuum assisted resin transfer method (VARTM). The composite was consolidated with the curing process: 90 °C for 2 h, 110 °C for 1 h, and 130 °C for 4 h in sequence, and the vacuum was about 0.1 MPa. As the single plain woven fabric ply is very thin, the 2-D plain woven composite was obtained by stacking 20 plies of fabrics along the thickness direction (in a 0° direction). The fiber volume fraction is 44.6% and 45.2% for 2.5-D angle-interlock woven composite and 2-D plain woven composite, respectively, obtained by muffle furnace combustion.

**Figure 1.** Structure of 2-D plain woven fabric and 2.5-D angle-interlock woven fabric. **Figure 1.** Structure of 2-D plain woven fabric and 2.5-D angle-interlock woven fabric.


**Table 3.** Manufacturer Specifications. **Table 3.** Manufacturer Specifications.

\* The density/(ends·cm<sup>−</sup><sup>1</sup> ) means the number of the yarns per centimeter. \* The density/(ends·cm−<sup>1</sup> ) means the number of the yarns per centimeter.

#### *2.3. Accelerated Ageing and Characterization 2.3. Accelerated Ageing and Characterization*

#### 2.3.1. Isothermal Ageing 2.3.1. Isothermal Ageing

The specimens were pretreated in an oven for 1 h at 80 °C and then divided into five groups. An unaged group (blank control group) and the other four groups were isothermally aged for 4, 8, 16, and 32 days at 180 °C in an air-circulating oven. Before ageing, all specimens were dried in the oven at 80 °C for 1 h. After ageing for a given time, the specimens were removed and cooled down to room temperature, then put into sealed bags to avoid moisture absorption. The specimens were pretreated in an oven for 1 h at 80 ◦C and then divided into five groups. An unaged group (blank control group) and the other four groups were isothermally aged for 4, 8, 16, and 32 days at 180 ◦C in an air-circulating oven. Before ageing, all specimens were dried in the oven at 80 ◦C for 1 h. After ageing for a given time, the specimens were removed and cooled down to room temperature, then put into sealed bags to avoid moisture absorption.

#### 2.3.2. Chemical Analysis 2.3.2. Chemical Analysis

Attenuated total reflectance Fourier transform infrared spectroscopy (ATR-FTIR) analyses (Nicolet 6700 FTIR spectrometer, ThermoFisher, Waltham, MA, USA) were used to determine the functional characteristics of neat resin in the surface layer before and after ageing for 32 days. Attenuated total reflectance Fourier transform infrared spectroscopy (ATR-FTIR) analyses (Nicolet 6700 FTIR spectrometer, ThermoFisher, Waltham, MA, USA) were used to determine the functional characteristics of neat resin in the surface layer before and after ageing for 32 days.

#### 2.3.3. Dynamic Mechanical Analysis (DMA) 2.3.3. Dynamic Mechanical Analysis (DMA) DMA (Q800, TA Instruments, New Castle, DE, USA) was performed on resin casting

*Polymers* **2022**, *14*, x FOR PEER REVIEW 4 of 18

DMA (Q800, TA Instruments, New Castle, DE, USA) was performed on resin casting in single cantilever mode with a frequency of 1 Hz over the temperature range 30 ◦C to 180 ◦C. The temperature ramping rate was 5 ◦C/min. in single cantilever mode with a frequency of 1 Hz over the temperature range 30 °C to 180 °C. The temperature ramping rate was 5 °C/min.

#### 2.3.4. Thermogravimetric Analysis (TGA) 2.3.4. Thermogravimetric Analysis (TGA) Thermogravimetric (TG) studies were carried out with a TA-Instrument (TGA 4000,

Thermogravimetric (TG) studies were carried out with a TA-Instrument (TGA 4000, PerkinElmer, Waltham, MA, USA). Powder samples of about 3 mg were heated under a nitrogen atmosphere. The samples were heated from 30 to 600 ◦C with a ramp rate of 20 ◦C/min. PerkinElmer, Waltham, MA, USA). Powder samples of about 3 mg were heated under a nitrogen atmosphere. The samples were heated from 30 to 600 °C with a ramp rate of 20 °C/min.

#### 2.3.5. Mechanical Test 2.3.5. Mechanical Test

The interlaminar shear tests of the specimens were conducted by the Instron universal testing machine (Instron 5967, Instron, Canton, MA, USA) at a test speed of 1 mm/min following ASTM D2344. The width-to-thickness ratio of the specimen is 2, while the span-to-thickness ratio is 4. The interlaminar shear tests of the specimens were conducted by the Instron universal testing machine (Instron 5967, Instron, Canton, MA, USA) at a test speed of 1 mm/min following ASTM D2344. The width-to-thickness ratio of the specimen is 2, while the spanto-thickness ratio is 4.

#### **3. Numerical Analysis 3. Numerical Analysis**

#### *3.1. Geometry Model 3.1. Geometry Model*

The mesoscale geometry models for both 2-D PWC (Figure 2a) and 2.5-D AWC (Figure 2b) were established based on measured structural parameters. The unit-cell of 2-D PWC consists of three fabric plies with random offset on account of the realistic staggered feature. The two models have a very close yarn volume fraction, which is 62.71% and 61.84% for PWC and AWC, respectively. The mesoscale geometry models for both 2-D PWC (Figure 2a) and 2.5-D AWC (Figure 2b) were established based on measured structural parameters. The unit-cell of 2-D PWC consists of three fabric plies with random offset on account of the realistic staggered feature. The two models have a very close yarn volume fraction, which is 62.71% and 61.84% for PWC and AWC, respectively.

Linear tetrahedral elements (C3D4) were chosen to generate the mesh of the yarns and matrix for both PWC (Figure 3a) and AWC (Figure 3b) models. Zero-thickness cohesive

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elements (COH3D6) were generated to represent the interface between the yarns and the matrix. The average mesh size of the model was set to be 0.2 mm. sive elements (COH3D6) were generated to represent the interface between the yarns and the matrix. The average mesh size of the model was set to be 0.2 mm.

Linear tetrahedral elements (C3D4) were chosen to generate the mesh of the yarns and matrix for both PWC (Figure 3a) and AWC (Figure 3b) models. Zero-thickness cohe-

**Figure 3.** Finite element models: (**a**) 2-D PWC; (**b**) 2.5-D AWC. **Figure 3.** Finite element models: (**a**) 2-D PWC; (**b**) 2.5-D AWC.

#### *3.2. Constitutive Model*

#### *3.2. Constitutive Model* 3.2.1. Epoxy Resin

3.2.1. Epoxy Resin The epoxy resin is treated as an elastic-plastic solid obeying the J2-isotropic hardening plasticity theory with an associated flow rule and a von Mises yield criterion. The properties of the epoxy matrix were experimentally characterized by a previous study [37] The epoxy resin is treated as an elastic-plastic solid obeying the J2-isotropic hardening plasticity theory with an associated flow rule and a von Mises yield criterion. The properties of the epoxy matrix were experimentally characterized by a previous study [37] and PWC, respectively.

#### and PWC, respectively. 3.2.2. Yarns

the yarns.

3.2.2. Yarns Woven yarns impregnated with the epoxy resin were regarded as transversely iso-Woven yarns impregnated with the epoxy resin were regarded as transversely isotropic unidirectional composite lamina. The compliance matrix of the yarns was obtained by the bridging model [38].

$$\mathbf{[S]} = \left(\mathbf{V\_{fy}} \left[\mathbf{S}^{\mathbf{f}}\right] + \mathbf{V\_{my}} [\mathbf{S}^{\mathbf{m}}] [\mathbf{A}]\right) \left(\mathbf{V\_{fy}} [\mathbf{I}] + \mathbf{V\_{my}} [\mathbf{A}]\right)^{-1} \tag{1}$$

**Carbon Fiber Epoxy Resin Fiber Tows**

−1

0.89 3.0

(1)

[S] = (Vfy[S f ] + Vmy[S <sup>m</sup>][A])(Vfy[I] + Vmy[A]) where [S f ] and [S <sup>m</sup>] refer to the compliance matrices of the fiber and resin. Vfy and Vmy are the volume fractions of the fiber and resin matrix in yarns, respectively. [A] is a bridging matrix and [I] is a unit matrix. A fiber packing fraction Vfy of 72% was determined by dividing the realistic fiber volume fraction of the specimen by the yarn volume fraction. Table 4 lists the basic properties of constituents and calculated engineering constants of where <sup>h</sup> S f i and [S <sup>m</sup>] refer to the compliance matrices of the fiber and resin. Vfy and Vmy are the volume fractions of the fiber and resin matrix in yarns, respectively. [A] is a bridging matrix and [I] is a unit matrix. A fiber packing fraction Vfy of 72% was determined by dividing the realistic fiber volume fraction of the specimen by the yarn volume fraction. Table 4 lists the basic properties of constituents and calculated engineering constants of the yarns.

<sup>E</sup>11(GPa) <sup>230</sup> 2.4 142.8 E<sup>22</sup> <sup>=</sup> E<sup>33</sup> (GPa) 14 6.4

*G23*(GPa) 5 2.3

**Table 4.** Basic properties of constituents and yarns.

G<sup>12</sup> <sup>=</sup> G<sup>13</sup> (GPa) 9


ν<sup>23</sup> 0.3 0.35

The yarn-matrix interfacial properties were described by a bilinear traction-separation constitutive model (Figure 4), relating the traction (t) and separation displacement (δ)

**Table 4.** Basic properties of constituents and yarns. = (2)

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#### 3.2.3. Interface evolution process. The interfacial stiffness degrades accordingly based on the damage var-

0

3.2.3. Interface

between two adjacent faces:

The yarn-matrix interfacial properties were described by a bilinear traction-separation constitutive model (Figure 4), relating the traction (t) and separation displacement (δ) between two adjacent faces: iable D varied between 0 (undamaged interface case) and 1 (complete decohesion case): = ( − 0 ) ( − 0 ) (4)

$$t = k\_p \tag{2}$$

where *t* is the nominal traction stress vector, consisting of three components: *tn*, *t<sup>s</sup>* , and *t<sup>t</sup>* , which represent one normal and two in-plane shear tractions respectively. *k<sup>p</sup>* is the penalty stiffness and *δ* is the vector of separations. where , refers to effective separation at the initiation of damage and complete failure, respectively, and refers to the maximum value of the effective separation attained during the loading history. If unloading occurs before the complete decohesion, the penalty stiffness will become ′ , where ′ = (1 − ).

**Figure 4.** Bilinear traction–separation constitutive model. **Figure 4.** Bilinear traction–separation constitutive model.

Table 5 listed the interfacial parameters for the model. Damage initiation is predicted by quadratic nominal stress criterion:

$$\left\{\frac{t\_n}{t\_n^0}\right\}^2 + \left\{\frac{t\_s}{t\_s^0}\right\}^2 + \left\{\frac{t\_l}{t\_l^0}\right\}^2 = 1\tag{3}$$

The Benzeggagh–Kenane (BK) fracture criterion [39] is adopted to control the failure evolution process. The interfacial stiffness degrades accordingly based on the damage variable D varied between 0 (undamaged interface case) and 1 (complete decohesion case):

$$D = \frac{\delta\_m^f \left(\delta\_m^{\text{max}} - \delta\_m^0\right)}{\delta\_m^{\text{max}} \left(\delta\_m^f - \delta\_m^0\right)}\tag{4}$$

where *δ* 0 *<sup>m</sup>*, *δ f <sup>m</sup>* refers to effective separation at the initiation of damage and complete failure, respectively, and *δ max <sup>m</sup>* refers to the maximum value of the effective separation attained during the loading history. If unloading occurs before the complete decohesion, the penalty stiffness will become *k* 0 *p* , where *k* 0 *<sup>p</sup>* = (1 − *D*)*kp*. *Polymers* **2022**, *14*, x FOR PEER REVIEW 7 of 18

Table 5 listed the interfacial parameters for the model.

**Table 5.** Interfacial parameters for cohesive layers [32]. **Table 5.** Interfacial parameters for cohesive layers [32].


3.2.3.1. Modeling of Matrix Shrinkage *3.3. Modeling of Matrix Shrinkage*

Matrix shrinkage leads to interface cracks, which are regarded as one of the primary thermal ageing damage modes in composites. The distribution of the shrinkage displacement field and the residual stress are closely related to the architecture of the composites. In this paper, the architecture effect was compared with the initiation and propagation of interface cracks in 2-D PWC and 2.5-D AWC. As the oxidation reaction process in terms of molecular dynamics is hard to reproduce using the finite element software ABAQUS, we employed the "shrinkage equivalent temperature difference method" [21,32] to simplify the shrinkage process of the matrix. The shrinkage displacement was reproduced by the temperature difference, which could cause the same deformation. The equivalent temperature difference can be defined by: Matrix shrinkage leads to interface cracks, which are regarded as one of the primary thermal ageing damage modes in composites. The distribution of the shrinkage displacement field and the residual stress are closely related to the architecture of the composites. In this paper, the architecture effect was compared with the initiation and propagation of interface cracks in 2-D PWC and 2.5-D AWC. As the oxidation reaction process in terms of molecular dynamics is hard to reproduce using the finite element software ABAQUS, we employed the "shrinkage equivalent temperature difference method" [21,32] to simplify the shrinkage process of the matrix. The shrinkage displacement was reproduced by the temperature difference, which could cause the same deformation. The equivalent temperature difference can be defined by:

$$
\Delta T\_{\rm sh}(t) = \frac{\varepsilon\_{\rm sh}(t)}{\alpha\_{\rm m}} \tag{5}
$$

(Stretch-

,

(in-plane deformation

where *εsh*(*t*) refers to matrix shrinkage strain at different ageing days; *α<sup>m</sup>* is the thermal expansion coefficient of the epoxy matrix; and ∆*Tsh*(*t*) is the equivalent temperature difference. where ℎ() refers to matrix shrinkage strain at different ageing days; is the thermal expansion coefficient of the epoxy matrix; and ∆ℎ () is the equivalent temperature difference.

The final resin shrinkage displacement was obtained from previous work [32] using the same material system. The temperature field load was applied to the elements within the oxidized layer using ABAQUS/Standard analysis. The boundary conditions are shown in Figure 5. The final resin shrinkage displacement was obtained from previous work [32] using the same material system. The temperature field load was applied to the elements within the oxidized layer using ABAQUS/Standard analysis. The boundary conditions are shown in Figure 5.

**Figure 5.** Boundary conditions for 2-D PWC and 2.5-D AWC. **Figure 5.** Boundary conditions for 2-D PWC and 2.5-D AWC.

#### **4. Results and Discussion 4. Results and Discussion**

*4.1. Thermal Oxidative Degradation Mechanisms of Epoxy Resin 4.1. Thermal Oxidative Degradation Mechanisms of Epoxy Resin*

#### 4.1.1. ATR-FTIR Analyses 4.1.1. ATR-FTIR Analyses

The FTIR spectra of neat resin obtained before and after ageing for 32 days at 180 °C are shown in Figure 6. Compared with the unaged sample, several distinct changes take place in the spectrum of the unaged resin: The characteristic band of C-H near 2800~3000 cm−<sup>1</sup> decreases in intensity. These phenomena demonstrate that C-H bonds were oxidized. The FTIR spectra of neat resin obtained before and after ageing for 32 days at 180 ◦C are shown in Figure 6. Compared with the unaged sample, several distinct changes take place in the spectrum of the unaged resin: The characteristic band of C-H near 2800~3000 cm−<sup>1</sup> decreases in intensity. These phenomena demonstrate that C-H bonds

ring structure was partly destroyed. The decrease of the absorption band near 1181 cm−1

which was derived from the bridge between the benzene rings [40], indicates that the backbone chains of epoxy resin can be cut by thermal aging. The bands at 1235 cm−<sup>1</sup>

Moreover, the characteristic absorption bands of the benzene ring near 1509 cm−<sup>1</sup>

ing vibration of -C = C- skeleton in the benzene ring) and 828 cm−<sup>1</sup>

were oxidized. Moreover, the characteristic absorption bands of the benzene ring near 1509 cm−<sup>1</sup> (Stretching vibration of -C = C- skeleton in the benzene ring) and 828 cm−<sup>1</sup> (in-plane deformation vibration of phenyl-H) decreased after 32 days of ageing, which shows that the benzene ring structure was partly destroyed. The decrease of the absorption band near 1181 cm−<sup>1</sup> , which was derived from the bridge between the benzene rings [40], indicates that the backbone chains of epoxy resin can be cut by thermal aging. The bands at 1235 cm−<sup>1</sup> attribute to aromatic ether increases after ageing. It can be inferred that aromatic ether is formed during the ageing process as an oxidation product [41]. *Polymers* **2022**, *14*, x FOR PEER REVIEW 8 of 18 attribute to aromatic ether increases after ageing. It can be inferred that aromatic ether is formed during the ageing process as an oxidation product [41].

**Figure 6.** IR spectra of the un-aged resin and the surface of an epoxy cube after ageing at 180 °C for 32 days. **Figure 6.** IR spectra of the un-aged resin and the surface of an epoxy cube after ageing at 180 ◦C for 32 days.

#### 4.1.2. Dynamic Thermomechanical Behaviors 4.1.2. Dynamic Thermomechanical Behaviors

Figure 7a shows the storage modulus over temperature after ageing at 180 °C. The storage modulus decreases gradually with the increase of ageing time in the glassy state due to chain scission. While in the rubbery state, the situation is reversed. The formation of the oxidized layer may account for this result. As mentioned above, chain scission occurs in the bridge between the benzene rings in DGEBA. Some small liberated segments escape from the system and molecular rearrangements may occur among the remainders, forming new compounds with a higher concentration of benzene ring, which have a higher glass transition temperature [16]. Figure 7a shows the storage modulus over temperature after ageing at 180 ◦C. The storage modulus decreases gradually with the increase of ageing time in the glassy state due to chain scission. While in the rubbery state, the situation is reversed. The formation of the oxidized layer may account for this result. As mentioned above, chain scission occurs in the bridge between the benzene rings in DGEBA. Some small liberated segments escape from the system and molecular rearrangements may occur among the remainders, forming new compounds with a higher concentration of benzene ring, which have a higher glass transition temperature [16].

0 d 4 d 8 d 16 d 0.4 0.6 0.8 (b)1.0 tanδ 0 d 4 d 8 d 16 d 32 d Figure 7b shows the loss factor (tan) over temperature. In this paper, the temperature at the peak value of tan δ is treated as the T<sup>g</sup> (glass transition temperature) of the sample. At the initial stage of ageing (0–8 d), the T<sup>g</sup> increases from 135 ◦C to 140 ◦C due to structural changes, such as further crosslinking and the loss of dangling chains, occurring slowly during the stage prior to the onset of severe degradation [42]. The structural changes also decrease the damping ability of the material, as indicated by the drop in the maximum tan δ values [43]. After 8 days, the ageing is dominated by thermolysis and the Tg decreases accordingly.

Figure 7b shows the loss factor (tan) over temperature. In this paper, the temperature at the peak value of tan δ is treated as the T<sup>g</sup> (glass transition temperature) of the sample.

20 40 60 80 100 120 140 160 180

Tem (℃)

0.0

0.2

20 40 60 80 100 120 140 160 180

Tem (℃)

32 d

0

500

1000

Storage Modulus (MPa)

(a)

1500

2000

2500

0.0

0.2

0.4

0.6

A

Absorbance

0.8

1.0

1.2

**Figure 7.** Dynamic thermomechanical behaviors. (**a**) storage modulus; (**b**) loss factor. **Figure 7.** Dynamic thermomechanical behaviors. (**a**) storage modulus; (**b**) loss factor. 4.1.3. Thermogravimetric Analysis

#### Figure 7b shows the loss factor (tan) over temperature. In this paper, the temperature 4.1.3. Thermogravimetric Analysis As the "skin-core" structure forms after thermo-oxidative ageing[44], the samples cut

at the peak value of tan δ is treated as the T<sup>g</sup> (glass transition temperature) of the sample. As the "skin-core" structure forms after thermo-oxidative ageing [44], the samples cut from the "skin" (i.e., the oxidized layer) and the inner core were tested, respectively (Figure 8). The temperature with 10% weight loss, T10%, is regarded as the initial decomposition temperature. The initial decomposition temperatures for the unaged sample and the oxidized layer and inner core are 383 ◦C, 374 ◦C, and 388 ◦C, respectively. Compared with the unaged one, the oxidized layer has a lower initial decomposition temperature, while the inner core's is higher. During the ageing process, some chain scission occurs in the sample by thermolysis and then the liberated segments migrate towards the surface of the sample [17]. Some liberated segments have poorer thermal stability, which leads to the advance of the initial decomposition temperature of the oxide layer. Meanwhile, the migration of small molecules towards the surface leads to the backward decomposition of the initial temperature at the inner core. Additionally, the weight remaining in the oxidized layer was about 16% at a temperature of 600 ◦C, which is much higher than that of the unaged sample. This indicates the formation of a more stable compound in the oxidized layer. from the "skin" (i.e., the oxidized layer) and the inner core were tested, respectively (Figure 8). The temperature with 10% weight loss, T10%, is regarded as the initial decomposition temperature. The initial decomposition temperatures for the unaged sample and the oxidized layer and inner core are 383 °C, 374 °C, and 388 °C, respectively. Compared with the unaged one, the oxidized layer has a lower initial decomposition temperature, while the inner core's is higher. During the ageing process, some chain scission occurs in the sample by thermolysis and then the liberated segments migrate towards the surface of the sample [17]. Some liberated segments have poorer thermal stability, which leads to the advance of the initial decomposition temperature of the oxide layer. Meanwhile, the migration of small molecules towards the surface leads to the backward decomposition of the initial temperature at the inner core. Additionally, the weight remaining in the oxidized layer was about 16% at a temperature of 600 °C, which is much higher than that of the unaged sample. This indicates the formation of a more stable compound in the oxidized layer.

The reinforcement architecture obviously affects the shrinkage displacement of the composite and further influences the distribution of ageing cracks [36]. In this section, the distributions of the matrix shrinkage and interface cracks in PWC and AWC were compared to reveal the effect of the reinforced structure on the thermal oxidative stability of

Figure 9 compares the local shrinkage of the matrix in PWC and AWC after thermal oxidative ageing for 16 days. The maximum calculated shrinkage depth was around 25 μm, which occurred in the resin rich zone of the top and bottom sides in both PWC and

attribute to aromatic ether increases after ageing. It can be inferred that aromatic ether is

unaged 32 d

**Figure 6.** IR spectra of the un-aged resin and the surface of an epoxy cube after ageing at 180 °C for

Figure 7a shows the storage modulus over temperature after ageing at 180 °C. The storage modulus decreases gradually with the increase of ageing time in the glassy state due to chain scission. While in the rubbery state, the situation is reversed. The formation of the oxidized layer may account for this result. As mentioned above, chain scission occurs in the bridge between the benzene rings in DGEBA. Some small liberated segments escape from the system and molecular rearrangements may occur among the remainders, forming new compounds with a higher concentration of benzene ring, which have a

formed during the ageing process as an oxidation product [41].

C

D

3000 2800 2600 2400 2200 2000 1800 1600 1400 1200 1000 800 600

)

4.1.2. Dynamic Thermomechanical Behaviors

higher glass transition temperature [16].

Wavenumber (cm**-**<sup>1</sup>

32 days.

B

**Figure 8.** Thermogravimetric analysis. **Figure 8.** Thermogravimetric analysis.

*4.2. Reinforced Structure Effect*

the composites.

AWC.

4.2.1. Matrix Shrinkage

## *4.2. Reinforced Structure Effect*

The reinforcement architecture obviously affects the shrinkage displacement of the composite and further influences the distribution of ageing cracks [36]. In this section, the distributions of the matrix shrinkage and interface cracks in PWC and AWC were compared to reveal the effect of the reinforced structure on the thermal oxidative stability of the composites.

### 4.2.1. Matrix Shrinkage

Figure 9 compares the local shrinkage of the matrix in PWC and AWC after thermal oxidative ageing for 16 days. The maximum calculated shrinkage depth was around 25 µm, which occurred in the resin rich zone of the top and bottom sides in both PWC and AWC. *Polymers* **2022**, *14*, x FOR PEER REVIEW 10 of 18

**Figure 9.** Local shrinkage displacement in PWC and AWC after thermal oxidative ageing for 16 days. **Figure 9.** Local shrinkage displacement in PWC and AWC after thermal oxidative ageing for 16 days.

As the shrinkage occurred, the boundary of the matrix still stuck to the yarn and internal stress was induced accordingly. The shrinkage displacement and stress induced by shrinkage were compared in two local areas between AWC and PWC (Figure 10). The AWC sample showed larger shrinkage displacement as well as internal stress compared to that of the PWC, which could be attributed to the different arrangement of the yarns. Taking six cross-sections from the surface to the interior (Figure 11), the yarn volume fraction on each surface varied from 45.4% to 77.8% in the PWC, and the fraction was kept at the constant value of 62.2% in the AWC. Meanwhile, the AWC has a larger yarn to yarn space than that of PWC, and the maximum shrinkage depth increased as the yarn to yarn As the shrinkage occurred, the boundary of the matrix still stuck to the yarn and internal stress was induced accordingly. The shrinkage displacement and stress induced by shrinkage were compared in two local areas between AWC and PWC (Figure 10). The AWC sample showed larger shrinkage displacement as well as internal stress compared to that of the PWC, which could be attributed to the different arrangement of the yarns. Taking six cross-sections from the surface to the interior (Figure 11), the yarn volume fraction on each surface varied from 45.4% to 77.8% in the PWC, and the fraction was kept at the constant value of 62.2% in the AWC. Meanwhile, the AWC has a larger yarn to yarn space than that of PWC, and the maximum shrinkage depth increased as the yarn to yarn spacing increased (Figure 12), as the extent of the matrix rich zone increased [29].

spacing increased (Figure 12), as the extent of the matrix rich zone increased [29].

**Figure 10.** Local shrinkage and internal stress in PWC and AWC with different ageing time. **Figure 10.** Local shrinkage and internal stress in PWC and AWC with different ageing time. **Figure 10.** Local shrinkage and internal stress in PWC and AWC with different ageing time.

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**Figure 12.** Comparison of shrinkage displacement with yarn distance in PWC and AWC. **Figure 12.** Comparison of shrinkage displacement with yarn distance in PWC and AWC.

#### 4.2.2. Interface Crack Evolution 4.2.2. Interface Crack Evolution 4.2.2. Interface Crack Evolution

Figure 13 shows that the evolution of interface cracks as ageing time increased in both AWC and PWC. The cracks were initiated near the ends of weft yarns and propagated along the warp yarns as the shrinkage-induced internal stress accumulated. With the increase in ageing time, the interface cracks propagated and linked together in the PWC, forming continuous interlaminar cracks. The cracks are prone to extending under the external loading, causing delamination of the PWC specimen [45–47]. Figure 13 shows that the evolution of interface cracks as ageing time increased in both AWC and PWC. The cracks were initiated near the ends of weft yarns and propagated along the warp yarns as the shrinkage-induced internal stress accumulated. With the increase in ageing time, the interface cracks propagated and linked together in the PWC, forming continuous interlaminar cracks. The cracks are prone to extending under the external loading, causing delamination of the PWC specimen [45–47]. Figure 13 shows that the evolution of interface cracks as ageing time increased in both AWC and PWC. The cracks were initiated near the ends of weft yarns and propagated along the warp yarns as the shrinkage-induced internal stress accumulated. With the increase in ageing time, the interface cracks propagated and linked together in the PWC, forming continuous interlaminar cracks. The cracks are prone to extending under the external loading, causing delamination of the PWC specimen [45–47].

**Figure 13.** Evolution of interface cracks with the increase of ageing time. **Figure 13.** Evolution of interface cracks with the increase of ageing time. **Figure 13.** Evolution of interface cracks with the increase of ageing time.

The cracks in the AWC were relatively discrete due to the existence of reinforced yarns toward the thickness direction, which restrained the connection of the cracks (Figure 14). Thus, the interlocked yarns along the thickness have a positive effect in improving the interlaminar properties of the AWC after thermal oxidative ageing. The cracks in the AWC were relatively discrete due to the existence of reinforced yarns toward the thickness direction, which restrained the connection of the cracks (Figure 14). Thus, the interlocked yarns along the thickness have a positive effect in improving the interlaminar properties of the AWC after thermal oxidative ageing. The cracks in the AWC were relatively discrete due to the existence of reinforced yarns toward the thickness direction, which restrained the connection of the cracks (Figure 14). Thus, the interlocked yarns along the thickness have a positive effect in improving the interlaminar properties of the AWC after thermal oxidative ageing.

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**Figure 14.** Distribution of interface cracks in AWC and PWC. **Figure 14.** Distribution of interface cracks in AWC and PWC. **Figure 14.** Distribution of interface cracks in AWC and PWC.

#### *4.3. Experimental Validates-Interlaminar Performance 4.3. Experimental Validates-Interlaminar Performance 4.3. Experimental Validates-Interlaminar Performance*

Figure 15a displays the typical load–displacement curves of AWC. The curves showed linear elastic in the initial stage before the first drop occurs after the peak load. After that, the load dropped in a "zigzag" manner due to the stress redistribution caused by local damage. In contrast to AWC, the curves of PWC (Figure 15b) showed a sudden drop in load after the peak load as failure took place catastrophically at the peak stress in a brittle manner. However, the brittleness decreased with the increase of the ageing time as indicated by the curves. Figure 15a displays the typical load–displacement curves of AWC. The curves showed linear elastic in the initial stage before the first drop occurs after the peak load. After that, the load dropped in a "zigzag" manner due to the stress redistribution caused by local damage. In contrast to AWC, the curves of PWC (Figure 15b) showed a sudden drop in load after the peak load as failure took place catastrophically at the peak stress in a brittle manner. However, the brittleness decreased with the increase of the ageing time as indicated by the curves. Figure 15a displays the typical load–displacement curves of AWC. The curves showed linear elastic in the initial stage before the first drop occurs after the peak load. After that, the load dropped in a "zigzag" manner due to the stress redistribution caused by local damage. In contrast to AWC, the curves of PWC (Figure 15b) showed a sudden drop in load after the peak load as failure took place catastrophically at the peak stress in a brittle manner. However, the brittleness decreased with the increase of the ageing time as indicated by the curves.

**Figure 15.** Load–displacement curves under the interlaminar shear tests: (**a**) AWC; (**b**) PWC. **Figure 15.** Load–displacement curves under the interlaminar shear tests: (**a**) AWC; (**b**) PWC. **Figure 15.** Load–displacement curves under the interlaminar shear tests: (**a**) AWC; (**b**) PWC.

Figure 16 compares the retention rate of the short beam strength along with ageing time. In the first eight days, the retention rate showed a slight declining tendency for both specimens, and the retention rate of the PWC was a bit higher than that of the AWC. While, with the increase of ageing time, the retention rate of PWC decreased significantly and became lower than that of the AWC after ageing for 16 days. Figure 16 compares the retention rate of the short beam strength along with ageing time. In the first eight days, the retention rate showed a slight declining tendency for both specimens, and the retention rate of the PWC was a bit higher than that of the AWC. While, with the increase of ageing time, the retention rate of PWC decreased significantly and became lower than that of the AWC after ageing for 16 days. Figure 16 compares the retention rate of the short beam strength along with ageing time. In the first eight days, the retention rate showed a slight declining tendency for both specimens, and the retention rate of the PWC was a bit higher than that of the AWC. While, with the increase of ageing time, the retention rate of PWC decreased significantly and became lower than that of the AWC after ageing for 16 days.

**Figure 16.** Retention rate of the short beam strength along with ageing time. **Figure 16.** Retention rate of the short beam strength along with ageing time. **Figure 16.** Retention rate of the short beam strength along with ageing time.

The obvious decline could be attributed to the change of failure modes after longterm exposure, as depicted in Figure 17. The obvious decline could be attributed to the change of failure modes after long-term exposure, as depicted in Figure 17. The obvious decline could be attributed to the change of failure modes after longterm exposure, as depicted in Figure 17.

**Figure 17.** Failure morphologies of AWC and PWC with different ageing time. **Figure 17.** Failure morphologies of AWC and PWC with different ageing time. **Figure 17.** Failure morphologies of AWC and PWC with different ageing time.

The failure modes of the AWC did not change much before and after ageing; the upper side of the specimen was under compression, leading to the initiation of interface

The failure modes of the AWC did not change much before and after ageing; the upper side of the specimen was under compression, leading to the initiation of interface

The failure modes of the AWC did not change much before and after ageing; the upper side of the specimen was under compression, leading to the initiation of interface cracks. The bottom side was under tension, and the yarn's breakage under the tension led to the final failure of the specimen.

The PWC shows different failure modes after long-term exposure. At the early stage of ageing (≤8 d), the failure is dominated by the fiber breakage at the bottom. At this stage, the PWC has good interface adhesion. The fiber tension failure occurred before the delamination failure. At this time, the short beam strength was determined by the tensile properties of the yarn, which cannot reflect the real interlaminar properties. The strength retention rate showed a slight decline consequently. With the increase of ageing time (≥16 d), the interlaminar cracks propagated, forming continuous interlaminar cracks. The cracks were prone to extending under the shear loading, causing the delamination failure of the specimen.

#### **5. Conclusions**

This paper investigates the thermo-oxidative stability of 2-D plain woven composites and 2.5-D angle-interlock woven composites using both experimental and numerical approaches. The specimens have been exposed to 180 ◦C in an air-circulating oven for 4, 8, 16, and 32 days. The combined effect of matrix degradation and interface cracks leads to the reduction of properties in polymer composites after ageing.

ATR-FTIR spectroscopy and thermal analysis were conducted to illustrate the degradation of the resin matrix during the ageing process. The combined effect of thermolysis and oxidation led to chain scission and molecular rearrangement. The newly formed oxidized layer had higher thermal stability compared to the inner core. The glass transition temperature of epoxy resin decreased due to thermolysis after long-term exposure.

The distribution of interface cracks is closely related to the reinforced structure of the composites, and a mesoscale finite element method has been established to illustrate the structure effect in terms of matrix shrinkage and crack evolution. The AWC sample showed larger shrinkage displacement as well as internal stress compared to that of the PWC due to its regular arrangement and a larger yarn to yarn space. Continuous interlaminar cracks were restrained in the AWC sample due to the existence of reinforced yarns toward the thickness direction, which had a positive effect on the improvement of the interlaminar properties of the material after thermal oxidative ageing.

Finally, the stability of interlaminar properties was experimentally estimated by short beam shear tests. The retention rate of interlaminar shear strength was 81.1% and 74.9% for AWC and PWC, respectively, after ageing for 32 days, and the results prove that the 2.5-D angle-interlock woven composites have better thermo-oxidative stability after long-term thermal exposure due to better structural integrity.

**Author Contributions:** Conceptualization, B.T.; Data curation, T.H.; Funding acquisition, M.C.; Investigation, T.H.; Methodology, B.T.; Supervision, J.Y.; Validation, J.Y.; Writing—original draft, M.C.; Writing—review & editing, X.G. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the National Natural Science Foundation of China (No. 12102144 and 52101385); the Zhejiang Provincial Natural Science Foundation of China (No. LY22E020013); the Jiaxing Public Welfare Technology Application Research Project (No. 2021AD-10012), the Open Project Program of Key Laboratory of Yarn Materials Forming and Composite Processing Technology of Zhejiang Province (No. MTC-2021-04 and MTC-2020-22); Natural Science basic Research Program of Shaanxi Province [grant numbers 2021JQ-659]; Research Fund for the Doctoral Program of Xi'an Polytechnic University [grant numbers 107020527].

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Conflicts of Interest:** The authors declare that they have no known competing financial interest or personal relationship that could have appeared to influence the work reported in this paper.

## **References**


## *Article* **Prediction of Long-Term Tensile Properties of Glass Fiber Reinforced Composites under Acid-Base and Salt Environments**

**Jihua Zhu <sup>1</sup> , Yangjian Deng <sup>1</sup> , Piyu Chen <sup>2</sup> , Gang Wang 3,\*, Hongguang Min 3,\* and Wujun Fang <sup>3</sup>**


**Abstract:** This study investigates the effects of deionized water, seawater, and solutions with various concentrations (5% and 10% by mass) of HCl and NaOH on the physical and mechanical properties of glass fiber reinforced polymers (GFRPs) through aging tests at 20 ◦C, 50 ◦C, and 80 ◦C. The tensile properties of GFRP were assessed by tensile testing at room temperature, and the strain during the tensile process was observed using digital image correlation. Additionally, the degradation mechanism was analyzed using scanning electron microscopy, and long-term tensile properties were predicted based on the Arrhenius model. The results indicated that the tensile strength of the GFRP decreased by 22%, 71%, and 87% after 56 d of exposure to 5% NaOH solutions at 20 ◦C, 50 ◦C, and 80 ◦C, respectively. The alkaline solutions had a more severe effect on the GFRP than deionized water, seawater, and acidic solutions. The experimental values and Arrhenius model predictions were found to be in good agreement with each other.

**Keywords:** GFRP; environmental degradation; tensile properties; DIC; long-term prediction

## **1. Introduction**

Pultruded glass fiber reinforced polymers (GFRPs), with high specific strength, nonmagnetic properties, and ease of forming, have the added advantage of economy over carbon and basalt fiber-reinforced composites, and they are widely used in aerospace, automotive and marine, medical, and civil engineering fields [1,2]. However, when compared with steel, pultruded GFRPs are less ductile and the failure mode of the material is usually brittle. When the GFRP is applied to critical structures, such as offshore drilling platforms, wastewater treatment plants, power plants, storage tanks, and pipelines in aggressive environments, it is often subjected to aggressive media such as seawater, high temperatures and humidity, as well as acidic and alkaline environments [3–6]. Under long-term effects of the aggressive media, the GFRP is prone to aging that leads to the degradation of the GFRP's structural performance and causes significant social and economic loss. To better promote the application of the GFRP, it is necessary to obtain performance degradation data for the GFRP in actual service and predict its long-term performance using accelerated life characterization methods with the aim of revealing the degradation mechanism of the GFRP and its degradation law. Studies have shown that water-based solution immersion is one of the conditions that causes the greatest degradation in GFRPs [7] because water absorption by GFRP causes hydrolysis and plasticization of the resin matrix, weakening the fiber-matrix interface [8,9].

Several studies have been conducted in recent years to investigate the durability of GFRP composites in highly corrosive environments. Shao and Kouadio [10] found that the

**Citation:** Zhu, J.; Deng, Y.; Chen, P.; Wang, G.; Min, H.; Fang, W. Prediction of Long-Term Tensile Properties of Glass Fiber Reinforced Composites under Acid-Base and Salt Environments. *Polymers* **2022**, *14*, 3031. https://doi.org/10.3390/ polym14153031

Academic Editors: Rushdan Ahmad Ilyas, Salit Mohd Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 4 July 2022 Accepted: 25 July 2022 Published: 26 July 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

polyester matrix GFRP did not show signs of water-absorption saturation after 260 d of immersion in water at 23 ◦C, despite a 17.5% decrease in tensile strength. At 70 ◦C, its water absorption increased and then decreased, even reaching negative values, with a 56.3% decrease in tensile strength; however, the Young's modulus of the GFRP remained almost unchanged at the two temperatures. Grammatikos et al. [11] reported similar findings and concluded that the composite increases moisture absorption and water diffusion owing to the increase in temperature, and the entire process of water diffusion does not reach saturation. Based on the Levenberg-Marquardt algorithm, Xin and Liu [12] concluded that there are three dominant factors for the change in mass due to moisture absorption: diffusion, polymer relaxation, and composite damage.

Sousa et al. [13,14] further compared the effects of salt solution and tap water on the deterioration properties of polyurethane GFRP and epoxy resin GFRP lap joints. They found that the residual strength of the GFRP joints was higher in salt water than in tap water, which is consistent with the test results [15,16]. Furthermore, a significant postcuring phenomenon was found; the maximum reduction in the ultimate load of the GFRP was 27% during the aging period of 730 d in the 20 ◦C water immersion tests, but at the end of the aging period, the ultimate load reached, or even slightly exceeded, the initial ultimate load. The post-curing phenomenon has also been identified in numerous studies [11,17–19] and is attributed to the fact that polymer systems containing styrene polyesters and vinyl esters undergo rapid local homo-polymerization during the curing process, but the composite does not achieve complete curing in a short time, which leads to further curing during the aging period.

Feng Peng et al. [18] comparatively studied the effects of a sulfuric acid solution with pH = 5 and a mass fraction of 30% on the epoxy resin GFRP at a temperature of 60 ◦C and found that a high-concentration sulfuric acid solution had severe corrosion effects on the GFRP. After 90 d of immersion, the bending properties of the GFRP decreased by 34% and 73%, respectively, and the hardness decreased by 6.5% and 55.6%, respectively. Kanerva et al. [20] studied the properties of vinyl ester in a sulfuric acid solution at a temperature of 90 ◦C, pressure of 1.5 MPa, and concentration of 50 g/l. They found that the bending stiffness of vinyl ester decreased by 22% after one year of acid leaching, but bending strength increased by 109%. Bazli et al. [21] studied and compared the degradation of the compressive properties of vinyl ester GFRP with different cross-sectional shapes under different erosion environments and found that the O-shaped cross-section was more prone to erosion than the I- and C-shaped cross-sections. Additionally, alkali was found to cause the most serious deterioration in GFRPs. After 147 d of exposure to an alkali solution with pH = 13.6 at 20 ◦C, compressive strength decreased by as much as 46%. Xue et al. [22] further found that the compressive strength of the GFRP decreased faster than the tensile strength in the alkali solution. In addition, Hota et al. [19] found that the interlaminar shear stress decay of a vinyl ester GFRP plate in an alkaline environment was equivalent to that of a sample cut from the plate, indicating that there was no size effect during degradation.

To further study the deterioration law of pultruded GFRP profiles in the acid-base and salt environments, the effects of different chemical media (deionized water, artificial seawater, hydrochloric acid solution, and sodium hydroxide solution of different concentrations) at 20 ◦C, 50 ◦C, and 80 ◦C on the deterioration performance of the pultruded GFRP profiles were experimentally studied. Subsequently, the quality and appearance changes, tensile properties, and micro morphology of the cross section of the GFRP were analyzed. The global tensile-strain field in the fracture process was analyzed using the digital image correlation (DIC) technique. Finally, the long-term tensile properties of the pultruded GFRP profiles were predicted using the Arrhenius model.

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

#### *2.1. Materials*

A GFRP plate pultruded by an epoxy resin and E-glass fiber produced by Mushi Composite Materials Technology (Suzhou) Co., Ltd. (Suzhou, China) was selected as

the raw material. The glass fiber was orthogonally braided and its mass fraction was 60%. To avoid stress concentration of the specimen during the tensile test, owing to the discontinuous notch formed by manual and mechanical cutting of the GFRP plate, the GFRP specimens used in the test were cut from the GFRP plate parallel to the fiber direction using laser-cutting technology. The specimen was cut into an I-shape, referring to *GB/T 1447-2005 Fiber-reinforced plastics composites determination of tensile properties*, with a nominal thickness of approximately 4 mm and length of 180 mm, as shown in Figure 1.

#### **Figure 1.** The specimen.

### *2.2. Aging Test*

The aging experiment explored and compared the degradation effects of four chemical media, namely deionized water (DW), artificial seawater (SW), sodium hydroxide solution (SH), and hydrochloric acid solution (HA), at different temperatures and ages on the GFRP in order to evaluate the law and mechanism of deterioration. In this test, acid and alkali solutions with mass fractions of 5% and 10% were used to further explore the degradation law of the GFRP under different concentrations of acid and alkali solutions. For the composition standard of artificial seawater, refer to *GB/T 3857-2017 Test method for chemical medium resistance of glass fiber reinforced thermosetting plastics*. The acid solution was prepared using concentrated hydrochloric acid at a concentration of 33% by mass, and the alkaline solution was prepared using analytically pure sodium hydroxide. Table 1 presents the test protocol for the deterioration tests. The GFRP specimens were soaked in six types of chemical media at three temperature gradients (20 ◦C, 50 ◦C, and 80 ◦C) with age settings of 28 d and 56 d. The specimens were numbered according to the following rules: the part before "-" indicates the immersion solution; 05 and 10 indicate the 5% and 10% mass percentage concentrations, respectively; the part after "-" indicates the immersion temperature. For example: 05SH-80, implies that the specimen was immersed in a 5% sodium hydroxide solution at an immersion temperature of 80 ◦C. The specimens were clamped and immersed in an acid-resistant, alkali-resistant, and high-temperature resistant airtight plastic box, as shown in Figure 2. Eighteen sets of test conditions were set, with 5 parallel specimens per set of conditions for a total of 180 specimens (=5 parallel specimens × 6 types of chemical media × 3 test temperatures × 2 ages). All specimens with an exposure temperature of 20 ◦C were placed in a cabinet at a temperature setting of 20 ± 0.5 ◦C, whereas the specimens with 50 ◦C and 80 ◦C conditions were placed in a water bath for heating.

**Figure 2.** GFRP specimen immersion test.

**Table 1.** Experimental scheme.


#### *2.3. Mass Changes*

The specimens were cleaned with running water, dried before and after the deterioration test, removed, and cooled to room temperature. The mass of the specimen was then measured (the mass of the specimen before and after exposure was noted as *m*<sup>0</sup> and *mt* , respectively, accurate to ±1 mg). The rate of mass change (*M*) was calculated using Equation (1).

$$M = \frac{m\_t - m\_0}{m\_0} \times 100\% \tag{1}$$

#### *2.4. Mechanical Testing*

After drawing auxiliary lines for each tested piece, a mechanical properties test was performed. The test apparatus were selected from the PWS-50 electro-hydraulic servo dynamic tester, and the tensile test at room temperature was performed at a constant speed of 0.2 mm/min with reference to *GB/T 1447-2005 Fiber-reinforced plastics composites determination of tensile* properties. Figures 3 and 4 show the test specimen and test setup with fixtures for the tensile test, respectively. Strain gauges were attached to both sides of the middle parallel section of each specimen parallel to the direction of tension and aligned centrally to obtain more accurate strain data in the initial stages of tension. An extensometer was used to indirectly measure strain over the entire scale.

**Figure 3.** Specimens.

**Figure 4.** Test setup.

## *2.5. Digital Image Correlation*

DIC is an optical metrology technique, based on digital image processing and numerical computation, which provides direct full-field displacement and strain with sub-pixel accuracy by comparing digital images of the test-object surface acquired before and after deformation [23]. To obtain a more visual and comprehensive understanding of the damage process in the GFRP during tension and deformation of specimens in various regions, one of the parallel specimens in each group was subjected to tensile testing using the DIC equipment. The DIC camera used was a Pointgrey-12.3 M, with a 4096 × 3000 resolution and a lens focal length of 50 mm; the acquisition interval was set to 400 ms. Before the test was performed, the specimens were marked with a scattering of black dots on a white background to allow the computer to better distinguish the movement of the scattering and to accurately analyze the displacement of the specimen.

### *2.6. Scanning Electron Microscopy*

A Phenom Pure scanning electron microscope (SEM) was used to obtain a reasonable explanation of the test phenomena from a microscopic perspective. As the subject of this test was a non-conductive material requiring a carbon or metal coating, the SEM specimen was first sprayed with a 9–12 nm platinum metal coating using sputter coating equipment and then placed in a sample cup to capture microscopic images.

#### **3. Results and Discussion**

#### *3.1. Appearance*

All specimens were bright green before the test; however, after the test, the appearance of the specimens changed, as shown in Figure 5. For the specimens tested at 20 ◦C, the appearance remained largely unchanged; for the specimens tested at 50 ◦C, the color faded from bright to light green; and for the specimens tested at 80 ◦C, the color degradation was more significant, with the surface resin matrix dissolving and essentially turning white. Additionally, the appearance of the specimens at different ages under the same conditions was slightly different, with the specimens showing a more pronounced color change at 56 d than at 28 d; the appearance of the specimens immersed in different chemical media under the same test conditions was also slightly different, with the specimens immersed in the sodium hydroxide solution showing the most pronounced change in appearance. The microstructural properties and chemical reactions of the samples are discussed further in this paper.


**Figure 5.** Appearance of specimens after immersion.

#### *3.2. Mass Changes*

Figure 6 shows the variation in the mass of the GFRP specimens over time for different hygrothermal environments. In general, mass absorption or loss under different degradation conditions increases with time; however, the rate of change can decrease or even tend to saturate. The change in mass is determined by a combination of the following: (i) water uptake, (ii) leaching of small molecular weight polymers, and (iii) leaching of dissolved hydrolysis products [4,5]. For specimens in the same chemical environment and exhibiting a consistent temperature effect, the degree of mass change increased with increasing immersion temperature. The masses increased by 0.09%, 0.34%, and 2.16% after 56 d of immersion in deionized water at 20 ◦C, 50 ◦C, and 80 ◦C, respectively. This was possibly because of the accelerated chemical reaction rate, owing to the increase in temperature, and a higher void pressure, resulting from the increased gas volume within the GFRP. This increase in pressure favors the extension of microcracks, thus increasing the free volume, within the specimen, that can be filled by the surrounding solution, such that the dendritic matrix absorbs more moisture [24].

In addition, it can be observed that the difference in mass change rates in the specimens, between the cases of deionized water and artificial seawater environments at 20 ◦C and 50 ◦C, is not significant; however, at 80 ◦C, the deionized water specimens absorb much more moisture than the artificial seawater environment specimens (the mass changes for DW-80-56 and SW-80-56 were 2.16 and 0.95%, respectively). This effect has been reported in earlier studies [13,25] and can be attributed to the cross-linking behavior of such polymers, which act as semi-permeable membranes. In the case of corrosive environments, such as acid and alkaline solutions, the specimens mainly showed mass loss (except for 05SH-80 and 10SH-80), contrary to the results [18,21]. This is because the dominant effect of mass change in the above environments is the decomposition of part of the polymer in the dendrimer matrix, its leaching into the solution [26], and dissolution of the resin on the

surface of the specimen [6]. In contrast, for specimens immersed in an alkaline environment at 80 ◦C, an increase in mass occurred (2.77% and 3.49% after 56 d in 05SH-80 and 10SH-80 solutions, respectively), probably owing to the increase in cavities within the laminate at higher temperatures resulting in increased water uptake [27]. The overall mass increase was observed when the weight-gain due to water absorption exceeded the mass loss due to leaching of the small molecular weight segment polymer, with the surface coating of the specimen being dissolved.

**Figure 6.** Change in specimen mass under various conditions.

#### *3.3. Tensile Properties*

3.3.1. Stress-Strain Curves and Failure Modes

Regardless of the exposure environment, all tested specimens showed greater degradation with increases in test temperature. As shown in Figure 7, the stress-strain relationship of various specimens exposed to 20 ◦C was similar, indicating that the elastic modulus of the tested GFRP specimen hardly decayed at 20 ◦C. With the increase in immersion temperature, the stress-strain curves of the tested specimens deviate from the control specimen and the maximum stress and strain decrease, indicating that the deterioration of the specimens is aggravated by the increase in temperature, especially for the specimen immersed in the 80 ◦C alkaline environment. Almost all the tested specimens exhibited failure near the clamping end (as shown in Figure 8a), which was mainly due to the stress concentration caused by the variable section near the clamping end of the specimen. However, the specimens immersed in acid at 50 ◦C and 80 ◦C exhibited delamination failure (Figure 8b). This is mainly because the binder resin is much worse than the fiber under high-temperature acid-leaching conditions, thus interlayer performance is weakened. Moreover, splitting occurs when the load reaches the limit [28].

#### 3.3.2. Tensile Strength and Modulus

The test data under all conditions are shown in Table 2. Figure 9 shows the average strength retention and error bars of the samples after 28 d and 56 d under different exposure conditions. The three exposure temperatures of 20 ◦C, 50 ◦C, and 80 ◦C are indicated by solid, scribed, and dotted lines, respectively, and different symbols are used for different exposure periods. Error bars indicate the magnitude of the difference between parallel samples when calculating the average strength.

**Figure 7.** Stress-strain curves of specimens at (**a**) 20 ◦C, (**b**) 50 ◦C, and (**c**) 80 ◦C.

**Figure 8.** Tensile failure modes: (**a**) fiber fracture, and (**b**) delamination.


**Table 2.** Summary of mechanical test results.

**Figure 9.** Variation of residual tensile strength as a function of immersion time and temperatures in various conditions.

As implied in Figure 9, regardless of the environment, temperature has a crucial impact on the deterioration of the specimens. Taking the sample immersed in artificial seawater as an example, after 56 d of exposure to 20 ◦C, the strength decreased slightly and the residual strength was 90.0%; however, when the temperature was raised to 50 ◦C, the deterioration was aggravated and the residual strength was 59.6%; at 80 ◦C, the degradation was further aggravated and the residual strength was only 40.8%. A hydrolysis reaction occurs when water comes in contact with the resin matrix. This reduces the compactness of the matrix, increases the number of pores, and reduces adhesion with the fibers, which may be the main reason for the worsening of the tensile properties of the GFRP [29]. The hydrolysis reaction between the resin matrix and water is shown in Equation (2). Hydrolysis of the resin polymer opens the ester bond and generates the corresponding carboxylic acid and free hydroxyl ions. An increase in temperature can accelerate the hydrolysis reaction, generate a higher void pressure conducive to microcrack propagation in the composite, and promote the entry of more water. However, it also affects the microstructure of the matrix, causing plasticization of the GFRP matrix [24,30] and weakening of the tensile properties of the GFRP. In addition, tensile strength does not decrease linearly with time, though most of the strength loss is generally observed after 28 d of exposure (except for DW-20), and the higher the exposure temperature, the more evident this phenomenon becomes. Taking 10SH-80 as an example, the residual tensile strength after 28 d immersion was only 24.1%, whereas the residual tensile strength after 56 d immersion was still 19.6%. It may be owing to the fact that the hydrolysis reaction of the composite is rapid, thus the rate of tensile strength loss is initially high before reducing due to the decrease in degradable reactants with the increase in degradation time. Consistent with the conclusions of many studies [7], the alkaline environment caused the most serious erosion of the GFRP. Specimens in 05SH-80-56 and 10SH-80-56 lost nearly all of their strengths, and the strength retention rates were only 12.6% and 19.6%, respectively. The tensile strength of the GFRP in deionized water and artificial seawater was lower than that of GFRP in acid. This may be because there are fewer free hydroxyl ions in the acidic environment, which inhibits the reaction between the glass fiber and hydroxyl ions to a certain extent, as shown in Equation (3). Sindhu et al. [30] suggested that fiber-matrix interface performance would be improved by immersion in an acid solution, thus improving tensile strength.

$$\text{R-COO-R'} + \text{H-OH} \rightarrow \text{R-COOH} + \text{R'} + \text{OH}^-\tag{2}$$

$$\text{-Si-O-Si-} + \text{OH}^- \rightarrow \text{Si-OH} + \text{Si-O-} \tag{3}$$

After 56 d of immersion in a deionized water environment at 20 ◦C, 50 ◦C, and 80 ◦C, the residual strengths were 86.2%, 53.3%, and 33.4%, respectively, compared to 90.0%, 55.4%, and 40.8%, respectively, for the seawater environment. Deionized water was found to reduce the tensile strength of the GFRP slightly more than artificial seawater. Compared to deionized water, the surface of the GFRP specimen is covered with a thin layer of salt in artificial seawater, which affects the exchange of matter between the interior and exterior of the specimen and, consequently, has less effect on the degradation of its tensile properties [13]. The effect of different concentrations of the sodium hydroxide solution and hydrochloric acid on the deterioration of the GFRP was also investigated. Unlike [31] and [32], it was found that a 5% sodium hydroxide solution deteriorated the GFRP more severely than the 10% solution, with the former having GFRP residual strengths of 77.9%, 29.3%, and 12.6% and the latter having the corresponding strengths of 81.7%, 30.6%, and 19.6% after 56 d of immersion at 20 ◦C, 50 ◦C, and 80 ◦C, respectively. One possible explanation is that, in the alkaline solution with higher concentrations of hydroxide, more hydrolysis products of the matrix and glass fibers were hoarded on the surface of the specimen at the beginning of the exposure, blocking the channels of the solution and inhibiting the matrix and glass fibers from undergoing hydrolysis reactions. In addition, it was observed that the strength degradation of the specimens immersed in 5% hydrochloric acid was slightly higher than the strength degradation of those immersed

in 10% hydrochloric acid, apparently because fewer hydroxide ions inhibited the degradation of the fibers, as shown in Equation (2).

Figure 10 shows the mean modulus of elasticity retention and error bars for the specimens after 28 d and 56 d under different conditions. The pattern of stiffness change was generally consistent with that of strength change. The 20 ◦C and 50 ◦C environments had less of an effect on the stiffness of the GFRP, with the remaining stiffnesses after 56 d of exposure for DW-50, SW-50, 05SH-50, 10SH-50, 05HA-50, and 10HA-50 being 89.9%, 87.2%, 92.5%, 92.0%, 88.3%, and 87.2%, respectively. However, when the exposure temperature was increased to 80 ◦C, the GFRP had severe deterioration, and the stiffness residuals after 56 d of immersion in the above environments were 81.5%, 82.0%, 56.0%, 57.2%, 73.1%, and 67.6%, respectively, which is very different from GFRP grids [33]. The Young's modulus of FRP composites is mainly determined by the modulus of elasticity of the fibers and resin matrix as well as the corresponding volume fraction [32], and it can be deduced that the glass fibers had significant deterioration at 80 ◦C. Overall, when comparing various exposure environments, alkali had the most significant effect on the stiffness of the GFRP, followed by acid, deionized water, and seawater.

**Figure 10.** Variation of the residual tensile elastic modulus as a function of immersion time and temperatures in various conditions.

#### *3.4. DIC Analysis*

Figure 11 shows the tensile stress field distribution and fracture location of the specimens after 56 d of exposure to artificial seawater, 5% hydrochloric acid, and a 5% hydroxide nano solution at 20 ◦C, 50 ◦C, and 80 ◦C. The first two figures from the left show the stress field distribution in the middle of the tensile test and when reaching the ultimate load, respectively, while the picture on the right side shows the fracture position of the specimen.

In Figure 11, the failure section of the specimens occurs at the maximum strain, which is in accordance with the maximum tensile strain theory. Furthermore, with the exception of specimens 05SH-50-56 and 05SH-80-56, the 'maximum strain' region for all specimens in the exposed condition was close to the clamping end on both sides, with fracture and failure occurring in those locations, which are consistent with Section 3.3.1. In contrast, the 05SH-50-56 and 05SH-80-56 specimens had failure sections in the middle parallel section because the specimens deteriorated to a greater extent in this environment than would be expected from the stress concentration phenomenon. Interestingly, the 'maximum strain' region in the specimen is not fixed at a certain place but may shift with the test process and eventually break at the maximum tensile strain, as shown in Figure 11a,f,i. This phenomenon is most evident in 05SH-50-56, where the specimen was the first to reach maximum strain and develop a crack near the lower fixture during tension; however, this was quickly followed by the development of a new crack above the previous, where it fractured completely.

Additionally, the strain fields of the specimens in different environments (with the exception of 05SH-50-56 and 05SH-80-56) were relatively smooth and uniform. This cor-

relates with the degree of material deterioration, and in relation to Section 3.3, the strain field for specimens with higher strength retention is smoother and more homogeneous. In contrast, the strain field distributions for specimens exposed to 05SH-50-56 and 05SH-80-56, which were the most severely weakened in terms of tensile properties, were patchy and discontinuous. This is due to the fact that pultruded GFRP profiles are generally relatively homogeneous. The GFRP which had not been severely eroded also showed relatively homogeneous deformation during tension. However, the distribution of defects within the GFRP, such as micropores, was not completely homogeneous, resulting in a considerable deterioration of the defective localities compared to other locations in an aggressive environment, which would be the first to induce greater strains and subsequently lead to a patchy distribution of the stress field.

**Figure 11.** *Cont*.

**Figure 11.** Tensile strain distribution exposed to SW, 5% NaOH, and 5% HCl for 56 d at 20 ◦C, 50 ◦C, and 80 ◦C, respectively. (**a**) SW-20-56, (**b**) 05HA-20-56, (**c**) 05SH-20-56, (**d**) SW-50-56, (**e**) 05HA-50-56, (**f**) 05SH-50-56, (**g**) SW-80-56, (**h**) 05HA-80-56, (**i**) 05SH-80-56.

#### *3.5. SEM Analysis*

Scanning electron micrographs of the fracture surfaces of the failed tensile specimens are shown in Figure 12. Comparing the microscopic morphology of the fracture surfaces after 56 d of immersion in SW, 05HA, and 05SH at 20 ◦C, 50 ◦C, and 80 ◦C, it was observed that there was no significant difference in the fracture morphology of the specimens exposed to different chemical environments at 20 ◦C. The fiber and resin sections were flat and compact, indicating that soaking at 20 ◦C had a good bonding effect between the fibers and resin, with no significant deterioration. When the ambient temperature was increased to 50 ◦C, gaps at the resin-glass fiber interface could be observed more clearly, with jagged fracture planes between the fibers. In particular, a large number of holes where the fibers were pulled out existed in the specimen section for the alkali environment. This indicates that, owing to the increased ambient temperature, the bonding between the fibers and resin matrix was substantially reduced, weakening the ability of the fibers to work perfectly with each other. When the immersion temperature was further increased to 80 ◦C, the bonding of the glass fibers to the epoxy resin in the specimen sections for all three chemical environments was further degraded (especially in the alkaline environment) and the nearly bare fibers, as well as the deboned strips of resin, were clearly visible, indicating that a large amount of resin matrix was withdrawn from the GFRP after 56 d of immersion at 80 ◦C. These observations support the discussions in Section 3.3.2 (Tensile strength and modulus).

**Figure 12.** SEM images of fractured sections of specimens. (**a**) SW-20-56, (**b**) 05HA-20-56, (**c**) 05SH-20-56, (**d**) SW-50-56, (**e**) 05HA-50-56, (**f**) 05SH-50-56, (**g**) SW-80-56, (**h**) 05HA-80-56, (**i**) 05SH-80-56.

Figure 13 illustrates the microscopic morphology of the peeled layer surface of a specimen exhibiting the delamination failure mode. Figure 13a,b clearly show the neatly defined grooves left by the peeled fibers, whereas Figure 13c shows the smooth fiber surface after the resin has been peeled. During the stretching process, the fibers on the surface of the sandwich that comprise the GFRP sheet peel off from the resin between the bonded sandwich, resulting in delamination. The delamination only occurred in specimens exposed to higher temperatures in acidic and seawater environments and at the later stage of the tensile test. A plausible explanation for this is that the resin bond between the interlayers of the specimens in these environments was severely degraded beyond the tensile capacity of the single interlayers, with splitting damage occurring before the specimens were damaged; for specimens in deionized water and alkaline environments, the degradation in tensile properties of the specimens was even more severe, with the tensile strength limit being reached before the specimens became delaminated. In addition, a comparison in Figure 13 shows that the higher the immersion temperature, the smoother the grooves left after the fibers are peeled, indicating more severe resin debonding. This explains the macroscopic

phenomenon in Section 3.3.1 that delamination damage is more likely to occur at 80 ◦C than at 50 ◦C in the same chemical environment.

**Figure 13.** SEM images of the specimens' strip surfaces. (**a**) 05HA-50-56, (**b**) SW-80-56, (**c**) 05HA-80-56.

## *3.6. Prediction of the Long-Term Behavior and Service Life of GFRP Tensile Strength*

Based on the results of the accelerated aging tests, the proposed deterioration model based on the Arrhenius theory [34] was used to predict the long-term tensile properties of the GFRPs. The following assumptions were made using the model: (1) Only one mode of chemical degradation can dominate the deterioration process, and the mode cannot vary with time or temperature. (2) The FRP must deteriorate in an aqueous solution but not in a dry environment [35]. The model considers temperature effects in accelerated aging tests and provides an accurate prediction for materials exposed to temperatures below their glass transition temperatures [36]. In this model, the relationship between performance retention Y and deterioration time t is given by Equation (4).

$$\mathbf{Y} = (100 - \mathbf{Y}\_{\infty}) \exp\left(-\frac{\mathbf{t}}{\tau}\right) + \mathbf{Y}\_{\infty} \tag{4}$$

where Y represents the retention of mechanical properties (%), t is the exposure time, τ is the regression fitting parameter, and Y<sup>∞</sup> is the retention of mechanical properties (%) for an infinitely long exposure time. Shenzhen, China, was chosen as the target site for the predictions.

In summary, the prediction procedure for each specimen in different chemical exposure environments consisted of four steps. Step 1 involved fitting the test results using Equation (4) to determine the regression fit parameters τ and Y<sup>∞</sup> at different temperatures. The fitting results are presented in Table 3.

Step 2 was for determining *<sup>E</sup><sup>a</sup> R* (i.e., the slope of the Arrhenius curve) of the samples for each exposure environment, where *E<sup>a</sup>* is the activation energy, and R is the universal gas constant. To this end, a linear fit of the logarithm (ln(t)) of the time taken for the sample retention (%) to reach a certain value (i.e., 95%, 90%, and 85% in this study) was plotted against 1000/K at different exposure temperatures (where K represents the absolute temperature). ln(t) at different temperatures can be determined from Equation (4) and the fitting parameters obtained in step 1. Figure 14 shows the Arrhenius plots and *<sup>E</sup><sup>a</sup> R* values for each sample type. The fitted straight lines for each type are nearly parallel, and the regression lines have R<sup>2</sup> values of at least 0.8106 (all above 0.80), indicating that the accelerated deterioration test is valid and that the model can be used to predict degradation in the tensile strength of the GFRP. The mean values of these slopes represent *<sup>E</sup><sup>a</sup> R* values, with different *<sup>E</sup><sup>a</sup> R* values indicating different degradation rates and possibly different degradation mechanisms.


**Table 3.** Regression coefficients in Equation (4) and time shift factors (TSF) for the GFRP specimens.

**Figure 14.** Arrhenius plots: specimens exposed to (**a**) DW, (**b**) SW, (**c**) 5% NaOH, (**d**) 10% NaOH, (**e**) 5% HCl, and (**f**) 10% HCl.

Step 3 was to calculate the time shift factor (TSF) for different exposure conditions using Equation (5).

$$\text{TSF} = \exp\left[\frac{E\_d}{R}\left(\frac{1}{T\_0} - \frac{1}{T\_1}\right)\right] \tag{5}$$

where *T*<sup>0</sup> is the lowest temperature (the annual average temperature in Shenzhen, China) and *T*<sup>1</sup> is the highest temperature (the exposure temperature in this study). Table 3 lists the TSFs of the six exposure environments.

In the final step, the TSF was multiplied by the corresponding exposure time at different temperatures to obtain the conversion time, and Equation (4) was used to plot the conversion time against the corresponding tensile strength retention to predict the long-term tensile performance of the GFRP, as shown in Figure 15. R<sup>2</sup> of the fitted line is at least 0.94, indicating that the Arrhenius model can accurately predict the deterioration pattern of the GFRP. The predicted results are in agreement with the experimental results, and it is evident that most of the performance loss occurs at early stages, regardless of the exposure environment. Tensile properties were predicted to decrease by almost 40% after only 89 d in a 5% sodium hydroxide environment at 22.3 ◦C, whereas GFRP showed the best resistance to weak-acidic environments, maintaining 49.4% tensile properties after 500 d in a 5% hydrochloric acid environment at 22.3 ◦C. Because the test environment is a long-time immersion in a solution, the above prediction results may be conservative compared to those of actual working conditions.

**Figure 15.** Predicted relationship between residual tensile strength and service life at the temperature of 22.3 ◦C for the GFRP exposed to various conditions.

#### **4. Conclusions**

The change in mass of GFRP increases with increasing ambient temperature and immersion time; however, its rate of mass change decreases with time gradually. In deionized water, artificial seawater, and alkaline solutions at 80 ◦C, the mass change of the GFRP was mainly characterized by an increase in mass; however, in acidic and alkaline solutions at 20 ◦C and 50 ◦C, it was mainly characterized by a mass loss.

The tensile strength and stiffness degradation of the GFRP were positively correlated with mass change. The residual tensile strength decreased with increasing ambient temperature and immersion time, but the rate of decrease gradually became slower with time. Compared to deionized water, artificial seawater, and acid, the alkaline environment had the greatest effect on the degradation of GFRP properties, with only 12.60% and 19.61% of the residual tensile strength after 56 d of exposure to 05SH-80 and 10SH-80, respectively, with degradation of tensile strength being more severe than that of tensile stiffness for GFRP.

The SEM results show that fiber-matrix debonding is most significant when the GFRP is exposed to alkaline environments, whereas the interlayer fibers of the GFRP are completely exposed in acidic and seawater environments at higher temperatures.

It can be deduced from the DIC image analysis that, at 50 ◦C and 80 ◦C, the GFRP corrodes uniformly when exposed to deionized water and acids but unevenly when exposed to alkaline solutions. The location of maximum strain may shift during tension, but failure nearly always occurs at the maximum strain.

The Arrhenius model was validated using experimental data, demonstrating that the model has good applicability. Compared to artificial seawater and alkaline environments, the GFRP is better at resisting long-term erosion in acidic environments. The tensile residual strength of the GFRP was predicted, using the Arrhenius model, to be 46.48% and 44.62% after exposure to 5% and 10% hydrochloric acid, respectively, for 2000 d in Shenzhen. The above test results and long-term performance prediction are of reference significance for the application of GFRP in Shenzhen.

**Author Contributions:** Conceptualization, all authors; methodology, J.Z., Y.D., P.C., G.W. and H.M; validation, Y.D. and P.C.; resources, J.Z.; data curation, Y.D. and P.C.; writing—original draft preparation, J.Z. and Y.D.; writing—review and editing, G.W., H.M. and P.C.; visualization, Y.D.; supervision, W.F.; project administration, W.F.; funding acquisition, G.W. and H.M. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was funded by the Key-Area Research and Development Program of Guangdong Province (grant number 2019B111107002) and the Major Projects of Central Research Institute of Building and Construction Co., Ltd., MCC Group (grant number JAA2018Kj01).

**Institutional Review Board Statement:** This study does not involve ethical research, so this statement is excluded.

**Informed Consent Statement:** Written informed consent has been obtained from the patient(s) to publish this paper.

**Data Availability Statement:** The data supporting this study are available from the corresponding author upon reasonable request.

**Acknowledgments:** We would like to recognize numerous co-workers, students, and research facility associates for giving specialized assistance on instrument use.

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

#### **References**


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