<figure><img src="image_236.png" /><figcaption>**FIGURE 5** | Fatigue crack growth rate data and linear fit used for Al 7075-T7531 tested at a load ratio of R = 0.1.</figcaption></figure>

this research. The entirety of the MMPDS data was used, as well as the upper portion (region II in the crack growth curve) of the Newman data. The data and fitted curve used in the computational study is shown in **Figure 5**. The fitted curve covered a range of stress intensity factors from 3.86 to 63.46 MPa $\sqrt{m}$, and yielded an *m* exponent of 3.93 and *C* equal to 9.04e-12 (m/cycle)/(MPa $\sqrt{m}$)$^m$.

# RESULTS

## Experimental Fatigue Crack Growth Behavior

Table 5 displays the fatigue test results for the AA7075 specimens cut from both middle thickness (t/2) and quarter thickness (t/4) of the rolled plate. In general, the material displayed a consistent number of cycles to failure regardless of the thickness at which the specimen was cut. However, 1 t/4 LCF specimen failed at 100 cycles. The likely reason for this significant deviation in fatigue life is related to the source of crack initiation and will be discussed in more detail in the discussion section. Based on this fatigue life data, it appears there is no preferred sampling thickness for this batch of AA7075-T7351.

The two HCF specimens exhibited almost identical crack length vs. cycle trends except with regards to crack initiation. The two LCF specimens also closely compare, where the dominant crack spent the majority of the fatigue life under a crack length of 250 µm. The t/2 and t/4 thickness of the plate displayed very minimal differences in both the crack growth rate and number of cycles to failure.

The crack growth rate trends for the LCF specimens are different than the trends for the HCF specimens. From the beginning of the crack growth for both LCF specimens, the crack growth rate was observed to fluctuate and is non-linear. This trend continued to the end of fatigue life.

## Fracture Surface Analysis

Through SEM, the source of crack initiation was determined to be a cracked intermetallic particle for each specimen observed in this

<table><caption>**TABLE 5** | Fatigue test results and particle size of crack initiating defects.</caption><thead><tr><td>Specimen</td><td>Stress range</td><td>Cycles to failure</td><td>Particle size (µm)</td></tr></thead><tbody><tr><td>LCF t/2</td><td>420</td><td>7,013</td><td>30.3</td></tr><tr><td>HCF t/2</td><td>260</td><td>55,365</td><td>18.4</td></tr><tr><td>LCF t/4</td><td>420</td><td>6,982</td><td>19.9</td></tr><tr><td>HCF t/4</td><td>260</td><td>58,308</td><td>14.7</td></tr></tbody></table>

study. In the vast majority of cases, a fatigue crack was found to nucleate from one or more particles at a single location and then propagate radially outward. In fact, we observed that a single specimen contained numerous large intermetallic particles (≈60 µm) throughout the fracture surface, but no noticeable fatigue crack was observed to initiate or propagate from these large particles. Of all the intermetallic particles that initiated fatigue cracks, the size ranged from approximately 10–30 µm. Upon scanning the intermetallic with EDS, all particles were determined to be iron-rich intermetallics. SEM images in backscatter electron (BSE) mode of representative fracture surfaces illustrating the initiation site for LCF and HCF regimens are shown in **Figures 6A–D**, respectively. A summary of these crack initiating particles based on SEM analysis is presented in **Table 5**.

EBSD was conducted on the surface of two specimens (one LCF t/2 and one LCF t/4) where the crack initiated and grain misorientation was measured. **Figure 7** displays the grain orientation map and misorientation map for the LCF t/2 (7(a)) and LCF t/4 (7(b)) samples. Large, equiaxed grains were observed based on EBSD scan results for the LCF t/2 specimen at the initiation site (Figure 7A). The scan on the t/4 specimen (Figure 7B) revealed an equiaxed grain structure with combination of large grains and small grains. High angle misorientation was observed between most of the identifiable grains for both the t/2 and t/4 LCF specimens. The dominant cracks observed in this study were observed to have initiated in a region with large grains and high grain misorientation.

The fatigue crack dimensions were measured on the fracture surface and used to determine the final crack aspect ratio and to compute the maximum crack length at the validity limit due to edge effects as indicated by the vertical lines in **Figure 8**.

## Fatigue Crack Growth Prediction

The results of the kinked model predictions compared to the experimental results are shown in **Figure 8** next to the geometries of the final crack configurations. The variation in the experimental crack growth rates is comparable to those determined in the FCGR characterization of this alloy (Jordon et al., 2012), where the variations in crack growth rate change at a analogous amplitude over a similar distance of crack growth as the crack grows through changing microstructure.

# DISCUSSION

## Microstructural Effects

The t/2 and t/4 thicknesses of the plate displayed very minimal differences in both crack growth rate and number of cycles to
<figure><img src="image_237.png" /><figcaption>**FIGURE 6** | Representative back-scatter electron (BSE) images of iron-rich intermetallic particles that initiated fatigue cracks in the (A), (B) low cycle fatigue (LCF) and the (C), (D) high cycle fatigue (HCF) regimes. Note that the iron-rich intermetallic particles are brighter compared to the darker aluminum matrix.</figcaption></figure>

failure, though the t/4 specimens consistently demonstrated a longer fatigue life to crack initiation, when compared to the t/2 specimens. Both EBSD scans conducted displayed a similar grain structure when comparing the t/2 and t/4 LCF specimens. The scans revealed an equiaxed grain structure with primarily large grains, and high grain misorientation was observed between most of the identifiable grains. Based on the data captured in this study, it appears there is no preferred plate thickness for fatigue loading. This is an important observation because similar observations on crack initiation mechanisms revealed that cracks can initiate from casting pores in thick rolled plates (Cauthen et al., 2020). From the results presented in Table 8, it is clear that neither thickness possessed preferred crack growth resistance and therefore differences in fatigue life are dictated by crack initiation and microstructure.

An abundance of iron-rich intermetallic particles was found on the fracture surfaces, where intermetallic particles were found to initiate the dominant crack. EBSD results revealed an equiaxed grain structure with primarily large grains and high grain misorientation at the location of surface crack initiation. Furthermore, no evidence of surface crack initiation from persistence slip bands (PSB) was observed in contrast to observations made in other high strength aluminum alloys with high grain misorientations (Cauthen et al., 2020).

However, due to the relatively high dispersion of the iron-rich intermetallic particles present in the legacy AA7075, it is not surprising that we observed cracks primarily initiating from these particles. These particles not only initiate cracks in AA7075 but can also influence the growth rates. The perturbation observed in small crack growth rates shown in Figure 13 has been shown to be influenced by discontinuities such as iron-rich intermetallic particles. In fact, crack growth rates have been observed to decrease significantly when cracks approached obstacles such as particles stringers in legacy AA7075. Once the driving forces exceeds the resistance provided by the iron-rich particle, the crack growth rate increases dramatically (Jordon et al., 2012).

## Model Performance

After accounting for the actual crack depth in the LCF t/2 specimen, as discussed below, crack growth modeling predicted the FCGR obtained from experimental results within the measurement error bounds and range of validity (Figure 8) for all four specimens. The model reasonably predicted the crack growth rate early in the crack life, even in areas known to have large microstructural effects that cause fluctuations in the growth rate and limit the validity of a linear elastic, SIF based approach.
<figure><img src="image_238.png" /></figure>

<figure><img src="image_239.png" /><figcaption>**FIGURE 7** | EBSD scans at the initiation site of the (A) LCF t/2 and (B) LCF t/4 specimens. Both sites were FIBed such that a smooth 50-100 µm wide surface remained for EBSD analysis.</figcaption></figure>

For all specimens tested, the model begins to severely over predict the growth rate at longer crack lengths. Note that this length scale is beyond the validity limit (shown as the vertical line in **Figure 8**) for which edge effects were estimated to be significant (greater than a 10%). This over prediction may be caused by plasticity induced crack closure effects, where crack closure occurs due to the combination of plastically deformed material left in the wake of the advancing crack and ahead of the crack tip (Walker et al., 2016).

The assumed crack depth also had a large effect on the predicted FCGR. The models were built on several assumptions, one of which was that the surface crack aspect ratio could be held constant at an average value with minimal error on the predicted FCGR. As seen in **Table 4**, as the crack becomes deeper, the surface geometry factor approaches one and converges to the SIF predicted by a through crack. For shallow cracks, the geometry factor is reduced. The geometry factor is applied to the effective SIF at each crack tip which is then input into the Paris law growth model, Eq. 9, where the crack growth is related to the SIF range raised to the *m* exponent. Therefore, the effect of the geometry factor error is raised by the power *m*. Thus, the effect of the error in the aspect ratio is higher for cracks with high SIFs than cracks with low SIFs, which helps explain the

increased error caused by the geometry factor in the LCF T/2 FCGR predictions.

For the LCF T/2 specimen, the crack was shallower than the assumed aspect ratio, causing the actual crack growth to be lower than the original prediction. The FCGR was recalculated using the final aspect ratio observed on the fracture surface of 0.6 for this specimen and the results are compared in **Figure 9**. The results using the actual aspect ratio show a much better fit when compared to the experimental results. To further explore the error caused by an incorrect aspect ratio, we investigated the results for the HCF T/4 specimen (actual aspect ratio of 1.08) for varying aspect ratios. FCGRs were calculated using the bounding aspects ratios of 2 (approaching a through crack) and of 0.5 (representing a shallow crack) as presented in **Figure 9**. The assumed crack depth is shown to have a large effect on the predicted growth rate, although the assumed value of 0.9 predicted the FGCR within the bounds of the experimental results. Therefore, it is concluded that it is necessary to use a realistic crack depth to determine the aspect ratio and subsequent geometry factor. If measurement error prohibits determination of a highly accurate crack aspect ratio, the problem can be bounded by a high and low estimation.
<figure><img src="image_240.png" /><figcaption>**FIGURE 8** | Model predictions vs. experimental crack growth for the (A) LCF t/2, (B) LCF t/4, (C) HCF t/4, and (D) HCF t/2. The red line shows the predicted value of crack growth at each crack length, and the green lines show the high and low estimates of experimental crack growth. The vertical blue line shows the rack length where significant edge effects begin.</figcaption></figure>

## Effects of Crack Shape

One purpose of this study was to investigate the usefulness of using DDT to model kinked surface fatigue crack growth. To do this, the SIFs predicted by the kinked crack model for the four specimens are compared to the SIFs predicted by modeling the cracks as straight. As shown in **Figure 10** for specimen LCF T/2 results, the higher fidelity gained by including the kinks in the crack shape had a relatively small, but noticeable effect on the predicted crack growth for the fatigue specimens investigated in this study.

A more substantial difference occurs in the mode II SIF between the straight and kinked models. The comparison for specimen LCF T/2 exhibited the largest divergence between

kinked and straight crack solutions, with more divergence between the solutions as the crack grows longer. This specimen configuration had a pronounced kink in the middle and displayed the fastest crack growth of the four specimens. As observed from a close inspection of the final crack configurations shown in **Figure 8**, the crack configurations were relatively flat or exhibited large kink opening angles, scenarios that are very close to a straight crack approximation and were loaded under tension applied perpendicular to the crack causing the importance of including kinks to be diminished.

To calculate the significance of including the kinks in the model prediction, a comparison between the experimental,
<figure><img src="image_241.png" /><figcaption>FIGURE 9 | (A) Effect of crack depth on the FGCR of specimen HCF T/4 for crack aspect ratios of 0.5, 0.9, and 2.0 (B) Experimental vs. model results for sample LCF T/2, showing model results assuming an aspect ratio of 0.9 vs. the actual aspect ratio of 0.6 (C) The fracture surface and final aspect ratio of the crack seen observed on sample LCF T/2.</figcaption></figure>

straight, and kinked model values was made via a comparison of the percent error, using Eqs 10 and 11.

$$ \text{Percent Error} = 100 * \frac{dc/dN_{\text{experimental}} - dc/dN_{\text{model}}}{dc/dN_{\text{experimental}}} $$

$$ \text{Percent improvement} = \% \text{ Error}_{\text{straight}} - \% \text{ Error}_{\text{kink}} $$

The percent improvement between the straight and kinked crack models was related to the amount of deviation from straight of the actual crack shape unsurprisingly. The straighter cracks

observed, in samples HCF T/4 and LCF T/4, showed smaller improvements between the models than the more kinked cracks, HCF T/2 and LCF T/2.

Within each sample though, differences between the performance of straight and kinked crack models appear to be primarily a function of difference between the angles, or directions, of the crack tips in each model. An example of this correlation is seen in the comparison of stress intensity factors between the straight and kinked models of sample LCF T/2 in Figure 10B. During the beginning of crack growth, when the observed crack was a sharp “V” shape, crack kinking of the crack tip angles of the real crack were captured better by the kinked model. Data points in this range had a higher percent improvement between the straight and kinked models, in the range of 30–52%, than the LCF T/2 crack as a whole. As the crack grew longer and became straighter as the crack tips turned toward a direction of maximum normal stress, the differences between the straight and kinked models decreased, even though the crack was modeled in more detail with multiple kinks. A similar observation can be seen in sample HCF T/2, when higher differences between the

<figure><img src="image_242.png" /><figcaption>FIGURE 10 | (A) Comparison of predicted FCGR for two LCF T/2 model configurations: a crack with a single kink and straight crack (B) Comparison of predicted FCGR (modes I, II, and effective) for the same two configurations.</figcaption></figure>
straight and kinked model are observed near the crack kinks, but then decreases as the crack grows outward in the direction of higher normal stresses away from the kink. These examples show that differences between a straight crack model and a kinked crack model are more severe when the kinking is closer to the crack tip, with the difference diminishing as the crack grows and straightens away from kinks.

## CONCLUSIONS

In this study, the nucleation and microstructurally small crack growth under fatigue loading for AA7075-T7351 was experimentally obtained, measured, and then modeled using a computational method based on DDT. In this work, microstructurally small fatigue crack growth was quantified using a novel surface replication process through interrupted fatigue testing. The replication process employed in this study used a fast-curing two-part epoxy that provided a submicron resolution of crack nucleation and propagation. This high-resolution surface replication process uniquely captured the microstructurally influenced fatigue crack growth behavior of the AA7075-T735 for the first time. In addition, using the spatial information taken from the replicas, the exact location for the fatigue initiation events on the fracture surface was known and subsequently imaged through scanning electron microscopy. As such, iron-rich intermetallic particles and clusters of particles were found to initiate fatigue cracks in this alloy. While iron-rich particles initiating fatigue cracks in AA7xxx alloys is not new information, the use the replica process coupled with post-mortem fracture surface analysis removed any ambiguity from the process of identifying sources of fatigue crack initiation.

Regarding the use of the surface replication process for model calibration, The high resolution replication technique provides critical input data for the prediction of the FCGR by providing a quantified crack shape. For the specimens investigated, crack growth modeling was found to predict experimental crack growth rates with an acceptable degree of accuracy. An important finding is that the geometry correction factor must be highly accurate, further emphasizing the importance of the high resolution technique to obtain crack shapes. Furthermore, the importance of explicitly modeling the kinks is exhibited when the angle or direction of the crack tips deviates from the shape of a straight crack, especially when the distance from the kink to the crack tip is small.

These results lead to the question of: what are the requirements in terms of crack configuration resolution for crack modeling to obtain converged and accurate results? Future work should address what amount of crack kinking/angle change is necessary and how many kinks/crack segments are required to obtain a given level of accuracy. Additionally, complex loading conditions may also necessitate higher fidelity modeling of the cracks. As a preliminary example, comparison of the results in Table 2 for biaxial loading demonstrate significant

differences in the SIFs between a kinked crack and a straight crack model. (The straight cracks in Table 2 were created by generating a straight line between the two crack tips of the corresponding kinked cracks).

Another important consideration is investigating the limits of prediction accuracy with regards to the geometry factor. This factor was successfully applied for this focused study on one specimen type using a factor derived for straight cracks. A comprehensive study is needed to determine the applicability of this assumption for cracks with more significant kinking and branching and to explore the need to vary the geometry factor per crack segment. An extensive sensitivity analysis should be performed to explore the effects of varying crack segment shape, crack segment length, crack depth relative to plate thickness, loading conditions, and number of cracks relative to the geometry factor on the FCGR prediction accuracy.

## Future Research and Development

The computational method was able to predict the FCGR for the cracks studied in this research. While the modeling resolution to include the effects of kink was not significant in this investigation, the computational foundation to predict the FCGR of kinked cracks was demonstrated. The method has been shown to be computationally efficient on a generic laptop computer and converges rapidly for a minimal number of DOF (measured by the number of opening terms included in an analysis). By integrating to obtain the solution, the method avoids many numerical approximations that increase the error in the results. Future applications of this method will benefit from the capability to obtain stress fields at discontinuities, including kinks and crack branching locations. Computational results also include full stress and displacement fields which will enable the prediction of crack path in future enhancements. Additionally, the method has been proven to converge for multiple crack arrays, even when the ligaments are very small between cracks (TerMaath et al., 2006). In summary, the method provides a fast yet accurate approach to evaluate existing crack patterns and conditions under general loading conditions and is applicable to small cracks through the use of the geometry factor. While the method is limited to modeling a crack array in a plate, this approximation does represent many real world scenarios. For example, this method has many potential applications for existing crack patterns that can be measured including:

* Informing inspection intervals

* Monitoring structural health

* Determining residual strength degradation

* Evaluating load-carrying capacity

* Informing repair and retirement strategies

* Locating areas where potential for new damage initiation is high (i.e., if corrosion, material defects, or overlooked cracks are present in a highly stressed area)

* Performing statistical studies using Monte Carlo simulation
DATA AVAILABILITY STATEMENT

The datasets presented in this article are not readily available
because Data was generated under projects with proprietary
restrictions by the sponsors. Requests to access the datasets
should be directed to steraaat@utk.edu.

AUTHOR CONTRIBUTIONS

RF – analytical framework development, calculations and results post-processing, sensitivity analysis and uncertainty quantification, paper writing. CC – experimental testing and results post-processing, analytical framework development, paper writing. SD – project coordination. JJ – project coordination, results post-processing, paper writing. AB – project administration. ST – analytical framework development and coding, calculations and results post-processing, paper writing, and project coordination.

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FUNDING

The authors would like to thank Boeing for the partial funding of
this project. Partial financial support was also provided by the
Office of Naval Research (ONR) Structural Reliability Program
through a Young Investigator Program grant (N00014-14-1-
0516/N00014-16-1-2370) under the direction of Dr. Paul Hess.

ACKNOWLEDGMENTS

Testing facilities used in this study were provided by the
University of Alabama and the Central Analytical Facility. The
authors thank the staff at the Central Analytical Facility for
assistance with microscopy. The authors also thank University
of Tennessee PhD students Cody Crusenberry and Corey Arndt
for their help with the UQ. The DICE software was provided and
developed at the Naval Surface Warfare Center Carderock
Division and the authors appreciate the technical support
from NAVSEA engineers Ken Nahshon and Nick Reynolds.

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**Conflict of Interest:** Author AB was employed by The Boeing Company.

The remaining authors declare that the research was conducted in the absence of
any commercial or financial relationships that could be construed as a potential
conflict of interest.

Copyright © 2021 Fleishel, Cauthen, Daniewicz, Baker, Jordon and TerMaath. This
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<figure><img src="image_243.png" /></figure>

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# Changing Compliance of Poly(Vinyl Alcohol) Tubular Scaffold for Vascular Graft Applications Through Modifying Interlayer Adhesion and Crosslinking Density

YeJin Jeong¹, Yuan Yao¹, Tizazu H. Mekonnen² and Evelyn K. F. Yim ¹,³,⁴*

¹Department of Chemical Engineering, University of Waterloo, Waterloo, ON, Canada, ²Institute of Polymer Research, Department of Chemical Engineering, University of Waterloo, ON, Canada, ³Waterloo Institute for Nanotechnology, University of Waterloo, Waterloo, ON, Canada, ⁴Center for Biotechnology and Bioengineering, University of Waterloo, Waterloo, ON, Canada

Poly(vinyl alcohol) (PVA) is a water-soluble polymer and forms a hydrogel that has been studied as a potential small-diameter (<6 mm) vascular graft implant. The PVA hydrogel crosslinked using sodium trimetaphosphate (STMP) has been shown to have many beneficial properties such as bioinert, low-thrombogenicity, and easy surface modification. Compared to conventional synthetic vascular graft materials, PVA has also shown to possess better mechanical properties; however, the compliance and other mechanical properties of PVA grafts are yet to be optimized to be comparable with native blood vessels. Mechanical compliance has been an important parameter to be studied for small-diameter vascular grafts, as compliance has been proposed to play an important role in intimal hyperplasia formation. PVA grafts are made using dip-casting a cylindrical mold into crosslinking solution. The number of dipping can be used to control the wall thickness of the resulting PVA grafts. In this study, we hypothesized that the number of dip layers, chemical and physical crosslinking, and interlayer adhesion strength could be important parameters in the fabrication process that would affect compliance. This work provides the relationship between the wall thickness, burst pressure, and compliance of PVA. Furthermore, our data showed that interlayer adhesion as well as chemical and physical crosslinking density can increase the compliance of PVA grafts.

**Keywords:** vascular graft, compliance, interlayer adhesion, crosslinking density, poly(vinyl alcohol), dip-casting

**OPEN ACCESS**

**Edited by:**
Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**
Guoping Guan,
Donghua University, China
Geoffrey Robert Mitchell,
Polytechnic Institute of Leiria, Portugal

**\*Correspondence:**
Evelyn K. F. Yim
eyim@uwaterloo.ca

**Specialty section:**
This article was submitted to Mechanics of Materials, a section of the journal Frontiers in Materials

**Received:** 15 August 2020
**Accepted:** 01 December 2020
**Published:** 14 January 2021

**Citation:**
Jeong Y, Yao Y, Mekonnen TH and Yim EKF (2021) Changing Compliance of Poly(Vinyl Alcohol) Tubular Scaffold for Vascular Graft Applications Through Modifying Interlayer Adhesion and Crosslinking Density. *Front. Mater.* 7:595295
doi: 10.3389/fmats.2020.595295

## INTRODUCTION

Small diameter synthetic vascular grafts, which are grafts with diameter <6 mm, are not commercially available due to having low patency. The low patency has been attributed to many mechanical and biological factors (Abbott et al., 1987). Among the different factors, mechanical compliance has been shown as one of the potential parameters that can influence the performance of small diameter vascular grafts (Jeong et al., 2020; Szafron, 2019). Compliance is defined in vascular graft engineering as the elasticity of the conduit when exposed to internal pressure (Szafron, 2019). The commercially available synthetic grafts, such as expanded polytetrafluoroethylene (ePTFE) and Dacron, have significantly lower compliance than native blood vessels (Tai et al., 2000). However, the
exact role compliance plays in the failure of small diameter vascular graft still remains elusive (Jeong et al., 2020). This is partially due to the difficulty in fabricating grafts using one material with variable, controllable, and predictable compliance.

One of the materials that is being studied as a potential off-the-shelf small-diameter vascular graft is poly(vinyl alcohol) (PVA) (Pohan et al., 2019; Cutiongco et al., 2015; Yao, 2020; Elshazly, 2004; Cutiongco et al., 2016b). PVA is a bioinert and low thrombogenic polymer that allows for easy chemical and topographical surface modification (Yao, 2020). In addition to having beneficial biomaterial properties, PVA has tunable mechanical properties (Cutiongco et al., 2016b). As a result, PVA is used in various applications. PVA vascular grafts are fabricated via dip-casting using cylindrical molds. PVA grafts have been shown to have dimensions and suture retention strength close to the native artery while withstanding burst pressure higher than that of the maximum systolic blood pressure (Chaouat, 2008). Although PVA grafts have higher compliance than ePTFE grafts, PVA grafts are still not as compliant as the native blood vessels (Pohan et al., 2020). Therefore, there is a need to investigate the different fabrication conditions of PVA grafts to further increase the compliance of PVA grafts. One limiting parameter to increasing compliance is the mechanical integrity of the grafts. As the grafts need to withstand the physiological pressure, increase in compliance cannot sacrifice the mechanical integrity of the grafts to ensure the safety of the patients. One of the parameters used to test the mechanical integrity of the vascular grafts is burst pressure. Burst pressure is the internal pressure at which the graft bursts, and is known to be closely correlated to the wall thickness of a tube (Menon, 2011). The motivation of this study is to make grafts with varying compliance using PVA, and identifying the burst pressure at respective compliance.

To further identify the parameters that can contribute to the mechanical properties of the PVA grafts, the interlayer adhesion and crosslinking density were explored. Interlayer adhesion and crosslinking density have been shown to contributed to dip-casting polymer and hydrogel mechanical properties. An increase in interlayer adhesion is known to increase the mechanical integrity of hydrogels (Kim et al., 2020). Interlayer adhesion is a critical factor since PVA grafts are fabricated through dip-casting multiple layers to reach the desired wall thickness. Additionally, the crosslinking density of the hydrogels can affect the mechanical properties of the polymer (Kim et al., 2020; Leone, 2010; Peppas et al., 2006). Increasing crosslinking density can result in stiffer hydrogels. PVA is a material with multiple mechanisms of chemical crosslinking with different chemical crosslinkers (Puy et al., 2019) and physical crosslinking with physical entanglement or via hydrogen-bonding (XIONG, 2008; Chen et al., 2016). In an effort to increase compliance while maintaining high burst pressure, interlayer adhesion and crosslinking densities were altered by altering fabrication process. We hypothesize that the increase in interfacial energy will allow for higher compliance and higher burst pressure when the grafts have the same wall thickness. Also, we hypothesize that a decrease in crosslinking density will result in higher compliance and lower burst pressure when the grafts have the same wall thickness. Lastly, we hypothesize that post-fabrication processing can cause physical crosslinking, and impact the compliance and burst pressure.

# MATERIALS AND METHODS

## Poly(vinyl alcohol) Vascular Graft Fabrication Process

Standard PVA crosslinking method was used for graft fabrication as described previously (Cutiongco et al., 2015). Briefly, cylindrical molds are dip-cast in the crosslinking solution, then dried in controlled environment as shown in **Figure 1A**. The complete list of the parameters that were varied during the fabrication of the grafts is shown in **Supplementary Table S1**. All of the grafts made using different fabrication parameters had 9 layers. The images of the grafts that were successfully fabricated using different fabrication parameters are shown in **Figure 1B**. The abbreviations and the fabrication conditions that was altered from the standard fabrication conditions is explained in **Figure 1C**.

## PVA Crosslinking Solution Preparation

For standard crosslinking density grafts, a 10% aqueous solution of PVA (Sigma-Aldrich, 85–124 kDa, 87–89% hydrolyzed) was mixed with 15% w/v sodium trimetaphosphate (STMP, Sigma-Aldrich) and 30% w/v NaOH at the volumetric ratio of 12:1:0.4, respectively, to create crosslinking solution. PVA grafts were fabricated using base-catalyzed STMP crosslinking. The PVA crosslinking solution was immediately dip-cast on a cylinder mold to form tubular PVA grafts.

To fabricate PVA grafts with different crosslinking density, 10% w/v PVA solutions of different molecular weights were used. Instead of the standard molecular weight of PVA (Sigma-Aldrich, 85–124 kDa, 87–89% hydrolyzed) low molecular weight PVA (Sigma-Aldrich, 13–23 kDa, 87–89% hydrolyzed) (LMW), or medium molecular weight PVA (Sigma-Aldrich, 31–50 kDa, 87–89% hydrolyzed) (MMW) were used while all other parameters were kept the same as the standard method. A low STMP (LS) concentration of 7.5% w/v was used to crosslink PVA with the standard PVA concentration and NaOH concentration. Lastly, low NaOH (LN) concentration of 15% w/v was used to fabricate grafts with decreased crosslinking density.

## Dip-Casting PVA Vascular Graft

PVA vascular grafts were cast as previously described (Figure 1A) (Pohan et al., 2019). In short, the molds for the tubular grafts were plasma cleaned, then dipped in PVA crosslinking solution. The coated molds were dried at 20–21°C and 30% humidity for varying waiting time, before dipping in the crosslinking solution again for the next layer. For the first four dips, the drying time between dips was 15 min, the drying time between dips in the fifth to seventh dip was 25 min, and drying time between dips for the eighth dip and higher was 30 min. The molds were inverted every-other dips to ensure even crosslinking throughout the mold. After the dipping completed, the PVA grafts were dried at 18°C and 70–80% humidity for 3 days.

Upon completion of the final drying step, the PVA grafts were rehydrated in 10x phosphate buffered saline (PBS) for 2 h, 1x phosphate buffered saline (PBS) for 2 h, then in deionized (DI) water for 2 h. For the samples used in phosphate quantification assay, samples were rehydrated in 9% w/v sodium chloride (NaCl) for 2 h, 0.9% w/v NaCl for
<figure><img src="image_244.png" /></figure>

C

<table><caption>**FIGURE 1** | Standard graft fabrication protocol and description of the varied parameter for each experimental groups. (A) Standard protocol, (B) representative images of the fabricated grafts for each experimental group, and (C) abbreviations and descriptions of the varied parameter for each experimental group. Samples for all groups were fabricated following the standard protocol except for the described parameter.</caption><thead><tr><th>Abbreviations</th><th>Description</th></tr></thead><tbody><tr><td>C</td><td>Control. Made following standard protocol (shown above). Samples had 9 layers.</td></tr><tr><td>LS</td><td>Low STMP. Made using 7.5% STMP instead of the standard concentration. Samples had 9 layers.</td></tr><tr><td>15W</td><td>15-minute waiting between dips. Made using 15-minute dry times between dips instead of the standard drying times. Samples had 9 layers.</td></tr><tr><td>30W</td><td>30-minute waiting between dips. Made using 30-minute dry times between dips instead of the standard drying times. Samples had 9 layers.</td></tr><tr><td>60-2D</td><td>Dried in 60°C post fabrication for 2 weeks. Made using the standard protocol, then dried at 60°C post fabrication for 2 weeks. Samples had 9 layers.</td></tr><tr><td>60-4D</td><td>Dried in 60°C post fabrication for 4 weeks. Made using the standard protocol, then dried at 60°C post fabrication for 4 weeks. Samples had 9 layers.</td></tr><tr><td>18-2D</td><td>Dried in 18°C post fabrication for 2 weeks. Made using the standard protocol, then dried at 18°C post fabrication for 2 weeks. Samples had 9 layers.</td></tr><tr><td>18-4D</td><td>Dried in 18°C post fabrication for 4 weeks. Made using the standard protocol, then dried at 18°C post fabrication for 4 weeks. Samples had 9 layers.</td></tr></tbody></table>

another 2 h, and then in DI water for 2 more hours. The grafts were removed from the molds once they were fully rehydrated. The completed grafts were then kept in DI water until experiments. This procedure was kept the same as the standard method for the different groups unless otherwise specified. Grafts with different drying times were also fabricated as a method to test how the dry time could vary the interfacial energy. The drying times were either consistent 15 min (15 W) or 30 min (30 W) between dips for these samples. In sample group to test the role of drying rate on

interfacial energy, PVA grafts were fabricated at higher humidity (HH) of 60 ± 5% during fabrication.

## Post-Fabrication Treatment of Extra-Dried PVA Vascular Grafts

After the rehydration step, some vascular grafts fabricated using the standard protocol were dried again to facilitate further crosslinking. Upon removal from the mold, the grafts were washed by submerging the grafts in DI water for 5 days. The grafts were dried again afterward at different conditions. The
post-fabrication drying conditions were: 2 weeks at 60°C (60-2D), 4 weeks at 60°C (60-4D), 2 weeks at 18°C (18-2D), and 4 weeks at 18°C (18-4D). Before testing, the grafts were rehydrated in DI water.

## Storage of the Fabricated Grafts

The samples were kept at room temperature at all times. All samples were stored in DI water in a 50 mL conical tube upon completion of the fabrication. DI water was poured out and replenished with fresh DI water every 7 days during storage.

## PVA Film Preparation

4 g ± 0.2 g of 10% w/v PVA solution was poured into petri dishes with diameter of 35 mm. The petri dish was covered and stored at 18°C chamber with a relative humidity of 70–80% for 3 days. The petri dish was then uncovered and stored at 18 C and 70–80% humidity for another 4 days. The dried PVA film (Film) was removed from the petri dish. Films were stored in petri dish until use.

## Inner-Diameter and Wall Thickness

2 mm segment rings were cut from the tubular PVA grafts at the ends and at the center of the samples. Images of the 2 mm segments were taken, and the inner diameter and wall thickness of the segments were measured for each segment from the images using MatLab. For each individual sample, the measurements of the inner diameter and the wall thickness from segments at different regions were averaged to yield the average inner diameter and the wall thickness of the sample.

## Compliance

3 cm segments of PVA grafts were exposed to 120 and 80 mmHg of hydrostatic pressure. Images of the grafts were taken at the defined pressures. The outer diameters of the grafts were measured, then the compliances of the grafts were calculated using the following equation (Chaouat, 2008):

$$ \frac{\text{diameter at 120 mmHg} - \text{diameter at 80 mmHg}}{\text{diameter at 120 mmHg}} \times 100 = \% \text{ compliance per 40 mmHg} \quad (1) $$

ePTFE vascular graft (Gore-Tex®, catalog number V04070L) with the inner diameter of 4 mm was used to measure the compliance of the ePTFE graft with comparable size using the stated setup.

## Burst Pressure

4 cm segments of PVA grafts were exposed to internal pressure induced by nitrogen gas at room temperature. The pressure was increased gradually until the graft burst. The burst pressure was recorded at the pressure at which the graft burst. A normalized burst pressure was calculated as burst pressure normalized per unit thickness by dividing the burst pressure of the samples by their wall thickness.

## Lap Shear Test

Lap shear test was performed following protocols described in published literature with modifications (Court et al., 2001). 1 cm

segments were cut from tubular PVA grafts. The segments were cut open longitudinally, then they were held in a flat position and dried at 60°C for 7 days. Segments prepared as such were partially rehydrated for 15 min and trimmed into 5 mm × 10 mm rectangular specimens. The surface of the specimens was patted dry using a paper towel. Then, the segment was glued onto two polylactic acid (PLA) holders so that the trimmed PVA segment was sandwiched between the two PLA holders. The PVA hydrogel segment was further rehydrated in DI water for 30 min. After rehydration, any residual water on the surface was removed. The non-overlapping parts of the holders were loaded onto a tensile testing equipment (AGSX Shimadzu, Japan), which was equipped with 1 kN load cell, and stretched at a rate of 25 mm/min until failure. Images of the samples after the test were taken using the laser confocal microscope (Olympus LEXT OLS5000-SAF, Japan) to measure the thickness variation within the destroyed samples. The heights of the samples were measured using Olympus data analysis software.

## Phosphate Quantification Assay

Phosphate quantification was performed according to the manufacturer protocol outlined in the phosphate quantification assay (Phosphate Assay Kit, Colorimetric, ab65622, Abcam). In brief, the rehydrated PVA grafts were dried at 60°C for 7 days. The dried grafts were then degraded in 10% nitric acid, then diluted to varying concentrations. The degraded polymer solutions were added into micro-plate and allowed to react with the assay reagent. The absorbances of the solutions were then measured. Using the absorbance, the amount of phosphate in the PVA graft was calculated.

## Fourier-Transform Infrared Spectroscopy

Fourier-transform infrared (FTIR) spectroscopy was performed following a published protocol (Pohan et al., 2020). PVA grafts were cut to 1 cm segments, and the segments were then cut open longitudinally. The grafts were dried flat. Fourier-transform infrared (FTIR) spectroscopy was performed using Thermo Fisher Scientific FTIR (Nicolett 6700) fitted with germanium and a high sensitivity pyroelectric detector. Percent transmittance of the samples were collected between 400 and 4,000 cm⁻¹. Sixty-four scans were acquired at a spectral resolution of 4 cm⁻¹. Data were normalized by modifying previously published protocol (Ahmed et al., 2019). The data was normalized to the highest % transmittance and the PVA film spectrum was subtracted from the % transmittance spectra.

## Differential Scanning Calorimetry

Differential scanning calorimetry (DSC; TA Instruments DSC Q2000) was performed following the published protocol (Pohan et al., 2020). In brief, the samples were dried at 60°C for at least 1 week 10 g ± 0.1 g of samples were measured and added into the Tzero Aluminum Hermetic pan. The heating cycle was set from 30°C to 260°C at the ramp rate of 10°C/min. Data were analyzed using the TA Universal Analysis software. Fractional crystallinity was calculated using heat of melting of 100% crystalline PVA of 138.60 J/g (Peppas and Merrill, 1976).
## Swelling Ratio

PVA grafts were submerged in DI water for at least a week to ensure thorough removal of charged ions from the grafts. Small sections of the tubular grafts were cut from the PVA grafts. Any residual DI water lingering on the PVA sections were removed prior to measuring the weights of the sections. The sections were then dried at 60°C for 5 days. The weights of the sections were measured every day for the subsequent days until the weight of the sections did not change further. Swelling ratio was calculated using the following equation (Lack, 2004):

$$ \text{Swelling ratio} = \frac{\text{weight of hydrated polymer sample}}{\text{weight of dried polymer sample}} \quad (2) $$

## Elastic Modulus

For longitudinal elastic modulus and ultimate strength measurement, the grafts were cut into 8 cm segments. The prepared samples were loaded onto the tensile machine (AGSX Shimadzu, Japan) with the gauge lengths of 2 cm. The samples were then stretched at a rate of 25mm/min until the tensile failure of the sample. For circumferential elastic modulus and ultimate strength measurement, the grafts were cut into 2 mm width rings. Curved holders were loaded onto the tensile testing machine grips. The PVA rings were loaded onto the holders so that the rings were held by the holders. The holders were then separated to achieve a gauge length of 5 mm. The samples were then stretched at a rate of 10 mm/min until the sample failed.

## Suture Retention Strength

Suture retention strength of the PVA grafts were measured using published protocol with modifications (Chaouat, 2008; Pohan et al., 2020). PVA grafts were cut into 4 cm segments. Using 7–0 suture, a single suture was made 5 mm from the edge of the segment for each sample. The remaining suture end was taped securely onto an immovable surface so that the sample was hanging by the suture. Load was applied at 10 g/min onto the sample until the sample broke. The weight at which the sample broke was measured.

## Statistical Analysis

Statistical significance was determined using a two-way ANOVA followed by Tukey's multiple comparison analysis with 95% confidence interval using MatLab. The data are presented as mean ± standard deviation. For all the presented data, *p* < 0.05 is denoted by * unless otherwise specified in the figure caption. The correlations between components were compared using principle component analysis (PCA). PCA plots were generated using JMP statistical software.

## RESULTS

### Compliance and Burst Pressure of the Grafts With Varying Wall Thickness

The relationships between the burst pressure vs wall thickness and compliance vs wall thickness are shown in Figure 2. The

cross-sectional images of the PVA grafts with different layers are shown in Figure 2A. The red arrow highlights the region in the graft that has different transparency. The wall thickness of the grafts increased linearly with the increased number of layers (Figure 2B). As shown in Figure 2C, a linear increase in burst pressure with increasing wall thickness was also observed. The compliance of the grafts displayed an inverse relationship with the wall thickness (Figure 2D).

### Properties of the Grafts Fabricated Using Different Fabrication Conditions

All of the grafts that were made using different fabrication parameters had 9 layers. The complete list of conditions used for fabrications and their fabrication outcomes are shown in Supplementary Table S1. The images of the grafts that fabricated successfully are shown in Figure 1B. The cross-sectional images of the successfully fabricated grafts are shown in Figure 3A. The wall thickness of the grafts was quantified (Figure 3B). LS, 15 W, 60-2D, and 60-4D had significantly thinner walls than the control group, while 30 W, 18-2D, and 18-4D showed no statistical difference from the control group. The wall thickness of 15 W was significantly thinner than 30 W. 60-2D and 60-4D did not show any statistical difference in wall thickness. Also, 18-2D and 18-4D did not show any statistical difference in wall thickness.

Despite having thinner walls than the control, grafts from LS, 60-2D, and 60-4D conditions had significantly lower compliance than the control group (*p* < 0.05, Figure 4A). Regardless of the post-fabrication drying temperature, the samples that were subjected to extra dehydration process post-fabrication had significantly lower compliance than the control group (Figure 4A). 30 W had lower compliance than the control group as well, despite having comparable wall thickness. Lastly, 15 W did not have statistically significant differences in compliance compared to the control group despite having thinner wall thickness. Within each group with the same fabrication or post-fabrication variations, all groups displayed an inverse relationship between wall thickness and compliance (Figure 4B). The burst pressure of 15 W was significantly lower than control, which was as expected due to the lower wall thickness of the group (Figure 4C). However, 60-2D and 60-4D had higher burst pressure than the controls despite having thinner walls. LS, 30 W, 18-2D, and 18-4D did not have significantly different burst pressure with respect to the control group. Within each group with the same fabrication or post-fabrication variations, all of the groups showed positive linear relationship between the wall thickness and the burst pressure (Figure 4D). Swelling ratio was used to estimate the crosslinking density. LS, 60-2D, and 60-4D had less swelling than the control group (Figure 4E). 15 W had higher swelling ratio than the control group. No correlation between the swelling ratio and the wall thickness was observed (Figure 4F). Further assessment of the burst pressure was performed by normalizing the burst pressure per unit thickness (Supplementary Figure S1). When the data was normalized to wall thickness, LS, 60-2D, and 60-4D had significantly higher normalized burst pressure than the control group.
<figure><img src="image_245.png" /></figure>

15 layer

<figure><img src="image_246.png" /></figure>

12 layer

<figure><img src="image_247.png" /></figure>

9 layer

<figure><img src="image_248.png" /></figure>

6 layer

<figure><img src="image_249.png" /></figure>

<figure><img src="image_250.png" /></figure>

Burst Pressure vs. Wall Thickness

<figure><img src="image_251.png" /></figure>

Compliance vs. Wall Thickness

<figure><img src="image_252.png" /></figure>

**FIGURE 2** | Physical and mechanical characterization of grafts with differential number of layers. (A) Cross-sectional images. Red arrows used to identify difference level of transparency between layers. (B) Wall thickness of the samples for differential number of layers. (C) Burst pressure plotted against the wall thickness. (D) Compliance plotted against the wall thickness. \*p < 0.05 between indicated groups. *n* = 9 for all groups. Error bar represents standard deviation.

Additionally, 60-2D and 60-4D had higher normalized burst pressure than 18-2D and 18-4D.

Phosphate content was quantified for the samples to verify the crosslinking density of the covalent crosslinking with STMP. There was no significant difference in phosphate content among the control, LS, 15W, 30W, 60-2D, and 18-2D (Figure 5A). On the other hand, both LS and 15 W had significantly higher phosphate content than the 18-2D and Film. FTIR was performed to identify if there were any changes in crosslinking between the groups or if there were any changes in functional groups introduced due to change in the fabrication process. Figure 5B shows FTIR data using normalized percent transmittance of the control group (C), LS, 60-2D, and 18-2D rehydrated in NaCl. All of the groups displayed peaks at similar wavenumbers, indicating there were no noticeable differences between samples in terms of functional groups. DSC was performed to further assess the degree of

crosslinking. **Figure 6** shows the DSC thermogram for control, 60-2D, and 18-2D. 18-2D had higher degree of crystallinity and crystallization temperature than the control group. Additional DSC data is shown in **Supplementary Figure S2**.

The interlayer adhesion was assessed in **Figure 7**. The shear strength of all of the groups were not significantly different from the control group. However, 15 W had the average shear strength of 3.75 MPa, which was more than twice higher than that of the control group (1.69 MPa). Statistical analysis between 15 W and control revealed *p* = 0.057. The interlayer adhesion of 15 W was also significantly higher than 30 W (*p* < 0.05). The 3D images of the polymers after the adhesive shear tests are shown in **Figure 5B**, then were used to acquire the height profile used to plot graphs in **Figures 7C-J**. These height profiles were then used to identify whether the polymers had high interlayer adhesion. 15 W had the highest slope, and therefore high interlayer adhesion, while other groups showed flat lines.
<figure><img src="image_253.png" /><figcaption>**FIGURE 3** | Physical characterization of the experimental groups. (A) Cross-sectional images. Red arrows used to identify difference level of transparency between layers within a sample. (B) Wall thickness of the samples for different number of layers. (C) Wall thickness of the samples. \*Indicates *p* < 0.05 with respect to control group. \# Indicates *p* < 0.05 between the groups being compared. *n* = 9 for all samples.</figcaption></figure>

**Supplementary Figure S3** shows graphs of the different properties measured plotted against either the wall thickness (**Supplementary Figures S1B-S3A**) or compliance (**Supplementary Figures S3C-S3F**). There was no obvious correlation between the measured properties with either the wall thickness or the compliance.

PCA analysis showed that the wall thickness and compliance showed a high positive correlation (Figure 8A). Also, it showed that the compliance and burst pressure showed a high negative correlation (Figure 8B). All other components did not display significant correlation (Supplementary Figure S4).

## Additional Analysis on Selected Groups

**Figure 9** shows the additional mechanical properties of selected groups. The groups with the highest and lowest compliance were chosen for the additional analysis. There was no statistical difference between all of the groups for all conditions except

for the LS for circumferential elastic modulus (Figures 9A-D). The suture retention strength of the groups also did not show any statistical difference (Figure 9E).

## DISCUSSION

Mechanical failure of vascular grafts can result in clinical catastrophe. To ensure the safety of the patients, vascular grafts were traditionally developed to have high mechanical properties matching the native blood vessel properties (Table 1 (Hansen et al., 1995; Fung, 1998; Kambric et al. 1984; Catanese III et al., 1999; Tai et al., 2000; Krejza et al., 2006; Sarkar, 2006; Konig et al., 2009; Markiewicz et al., 2017; Wise, 2011). Compliance, the elasticity of the blood vessel when exposed to internal pressure, is one of the important parameters that is being studied in vascular graft engineering to understand
<figure><img src="image_254.png" /><figcaption>**FIGURE 4** | Bulk properties of the samples. (A) Compliance of each group. (B) Compliance of each group plotted against their respective wall thickness. Average values of compliance and wall thickness of the internal replicates used to plot the graph. (C) Burst pressure of each group. (D) Burst pressure of each group plotted against their respective wall thickness. Average values of burst pressure and wall thickness of the internal replicates used to plot the graph. (E) Swelling ratio of each group, and (F) swelling ratio plotted with respect to the wall thickness of each group. *Indicates *p* < 0.05 with respect to control group. # Indicates *p* < 0.05 between the groups being compared. *n* = 9 for all samples.</figcaption></figure>

the current limitations of small diameter vascular grafts (Szafron, 2019; Jeong et al., 2020). There are methods to increase the compliance of a material. One of the methods is to reduce the wall

thickness (Biglino, 2013; Nezarati et al., 2015; Gao, 2019). However, reducing the wall thickness can negatively affect the burst pressure (Menon, 2011; Zhu and Leis, 2012). Using vascular

<table><caption>**TABLE 1** | Reported values of mechanical properties of native blood vessels and expanded polytetrafluoroethylene (ePTFE).</caption><thead><tr><th></th><th>Diameter (mm)</th><th>Burst pressure (mmHg)</th><th>Compliance (% per 100 mmHg)</th><th>Elastic modulus (MPa)</th></tr></thead><tbody><tr><td>Internal mammary artery (human)</td><td>1 – 2 (Markiewicz et al., 2017)</td><td>2,267 ± 215 (Wise, 2011)</td><td>12 (Konig et al., 2009)</td><td>0.267 ± 0.046 (Wise, 2011)</td></tr><tr><td>Carotid artery (human)</td><td>5.11 ± 0.87 (Krejza et al., 2006)</td><td>5,000 (Sarkar, 2006)</td><td>6.6 ± 1.3 (Hansen et al., 1995)</td><td>7 – 11 (Fung, 1998)</td></tr><tr><td>Saphenous vein (human)</td><td>4.9 ± 0.2 (Tai et al., 2000)</td><td>2,295 ± 272 (Wise, 2011)</td><td>5.0 ± 6.7 (Tai et al., 2000)</td><td>0.683 ± 0.147 (Wise, 2011)</td></tr><tr><td>ePTFE</td><td>5.4 ± 0.1 (Tai et al., 2000)</td><td>2,590 – 3,626 (Kambric et al., 1984)</td><td>1.2 ± 0.3 (Tai et al., 2000)</td><td>42.9 ± 2.6 (Catanese III et al., 1999)</td></tr></tbody></table>
<figure><img src="image_255.png" /></figure>

**FIGURE 5** | Functional groups and STMP crosslinking density in PVA films. (A) Quantification of phosphate content for the samples to characterize covalent crosslinking by STMP. *Indicates *p* < 0.05 between the indicated groups. *n* = 6 for all groups. C PBS was control sample rehydrated in NaCl solution. *Indicates *p* < 0.05 with respect to the control groups. #Indicates *p* < 0.05 between the groups being compared. (B) Fourier-transform infrared spectroscopy data. Data was normalized to PVA films. *n* = 3 for all samples.

grafts with low burst pressure could become a critical problem for the patients with the vascular graft implants. Therefore, it is important to understand the effects different parameters have on the burst pressure when increasing the compliance of the vascular grafts. Two parameters that have been found to affect the mechanical properties of hydrogels are interlayer adhesion and crosslinking density (Leone, 2010).

Crosslinking of PVA has been extensively studied. A well-established physical crosslinking method of PVA is freeze-thawing. PVA grafts with freeze-thawing crosslinking are mainly fabricated by bulk molding fabrication. PVA hydrogels crosslinked using freeze-thawing have been shown to have tailored mechanical properties by varying the number of freeze-thawing cycles (Peppas, 1975; Gupta et al., 2012; Holloway et al., 2013). The freeze-thawing crosslinking of PVA involves phase separation and crystalline formation. Hydrogen bonds form in the aqueous solution of PVA. During the freezing of PVA solution, ice forms. The formation of ice induces the growth of polymer crystallinity. The

growth of polymer crystallinity acts as physical crosslinking points. Subsequently, water-insoluble hydrogels are formed in the later thermal cycles. The abundant hydroxyl groups in the backbones of PVA polymer chains also provide reactive sites for hydrogen bonding for chemical crosslinking of PVA (Otsuka and Suzuki, 2009). One of the most commonly used crosslinking agent for PVA is glutaraldehyde. It reacts with hydroxyl groups to form an ester bond. Chaouat 2008) found that the hydroxyl groups on PVA can also react with STMP to crosslink PVA hydrogels by phosphoesterification. The PVA hydrogels chemically crosslinked with STMP were shown to be capable of achieving a wide range of mechanical properties by altering the fabrication parameters such as humidity and water evaporation (Cutiongco et al., 2016b). A recent study comparing the hemocompatibility of PVA hydrogels crosslinked by STMP, glutaraldehyde, and freeze-thawing found that crosslinking methods also affect the biological performance of PVA (Bates et al., 2020).
<figure><img src="image_256.png" /><figcaption>**FIGURE 6** | Degree or crosslinking. Differential Scanning Calorimetry (DSC) data. (A) DSC of (i) control, (ii) 60-2D, and (iii) 18-2D. Data is shown for the second heating and cooling cycle. (B) Heat of fusion and fractional crystallinity of the samples. n = 1 for all groups.</figcaption></figure>

PVA vascular grafts are fabricated using dip-casting with cylindrical molds (Cutiongco et al., 2016a). Different mechanical properties of the PVA vascular grafts can be achieved by altering the number of layers and fabrication parameters. We have observed heterogeneity in transparency between the layers (Figure 2A, red arrows) in PVA vascular grafts. The observed heterogeneity can be the result of inconsistencies in crosslinking between the layers (Peacock and Calhoun, 2012). The heterogeneity can also be an indication of a weak interlayer adhesion. A potential method of overcoming the weak interlayer adhesion would be increasing the interlayer crosslinking (Yu et al., 2013). With PVA hydrogels, the interlayer crosslinking could occur by having hydrogen bonding between the residual -OH groups of the PVA layers. It was hypothesized that through increasing the interlayer adhesion, the resulting grafts can have high compliance while maintaining burst pressure. Moreover, the heterogeneity suggests that the crosslinking density can be further fine-tuned. We hypothesized that a decrease in the chemical crosslinking density will result in higher compliance while causing a decrease in burst pressure. Furthermore, we hypothesize that post-fabrication processing can affect physical crosslinking, and impact the compliance and burst pressure.

## Compliance and Burst Pressure of the Grafts With Varying Wall Thickness

Most often, tensile measurements are made to assess the mechanical properties of the polymers (Abraham et al., 1998;

Peiffer and Nielsen, 1979). Additional properties, such as residual stress, rheology, Izod impact measurement, etc. (Peiffer and Nielsen, 1979; Abraham et al., 1998), are also tested depending on the application of the polymers. For vascular grafts, burst pressure and compliance are important parameters to measure. Although PVA vascular grafts have been studied previously (Elshazly, 2004; Chaouat, 2008), the relationship between the wall thickness and burst pressure and compliance have not been characterized. Thorough characterization of the compliance and burst pressure of the grafts with different wall thickness was performed first to identify the compliance and burst pressure achieved by different wall thickness. The relationship between the wall thickness and burst pressure (Figure 2C) was expected as the structural integrity of tubes is dependent on the wall thickness (Menon, 2011; Zhu and Leis, 2012). The grafts with 6 layers displayed high compliance, but its burst pressure was only 232.7 ± 22.8 mmHg. As systolic blood pressure can rise to 220 mmHg during exercises, PVA grafts with 6 layers were deemed too weak. Grafts with 9 layers have a burst pressure of 485.5 ± 17.9 mmHg, which is at least double the systolic pressure of 220 mmHg. However, the compliance of the grafts with 9 layers was 2.6 ± 0.54% per 40 mmHg. Furthermore, compliance of the PVA grafts with more than 9 layers had compliance <2% per 40 mmHg. As shown in Table 1 (Kambric et al. 1984; Hansen et al., 1995; Fung, 1998; Catanese III et al., 1999; Tai et al., 2000; Krejza et al., 2006; Sarkar, 2006; Konig et al., 2009; Wise, 2011; Markiewicz et al., 2017) and Figure 4A, the compliance of the PVA grafts with 9 layers and ePTFE are lower than the native
<figure><img src="image_257.png" /><figcaption>**FIGURE 7** | Interfacial energy. (A) The adhesive strength of groups. 15 W had *p* = 0.0571 with respect to the control group, and *p* = 0.0495 with respect to the 30 W. (B) Height map of the samples after adhesive shear test. (C)–(J) Height profile of the samples. \*Indicates *p* < 0.05 with respect to control group. \# Indicates *p* < 0.05 between the groups being compared. *n* = 3 for all samples.</figcaption></figure>

blood vessels. After assessing the compliance and burst pressure of the groups, the PVA vascular grafts with 9 layers was chosen for further investigation as they had the highest compliance with good burst pressure. Using the PVA grafts with 9 layers, the effect of crosslinking density and interlayer adhesion on the compliance and burst pressure were assessed.

## PVA Grafts Fabricated Using Different Fabrication Conditions

The standard fabrication method is described in **Figure 1A** in detail. The crosslinking of PVA using crosslinker STMP in NaOH solution has been described in the literature numerous times

(Elshazly, 2004; Chaouat, 2008; Leone, 2010; Ino et al., 2013). In brief, the -OH groups on PVA loses the hydrogen atom in the highly basic solution created by NaOH, and becomes reactive. This reactive site then reacts with the STMP. Subsequently, another PVA molecule reacts with the STMP and results in a crosslinked polymer. This process can be modified via changing the PVA concentration, STMP concentration, NaOH concentration, temperature, humidity, and waiting time between the dips (Chaouat, 2008). The fabrication conditions were varied to achieve different level of interlayer adhesion and crosslinking density.

The interlayer adhesion occurs when there are crosslinks that spans multiple layers of polymers. The mechanical properties of PVA vascular grafts can be affected by the interlayer adhesion as
<figure><img src="image_258.png" /><figcaption>**FIGURE 8** | Principle component analysis (PCA) for the identification of significant components. (A) PCA plot of component 1 vs. component 3. (B) PCA plot of component 2 vs. component 3.</figcaption></figure>

PVA graft fabrication method is a dip-casting process with waiting periods between dips. The dip-casting process is a type of polymer fabrication method where the mold is dipped multiple times to achieve the desired thickness (Abraham et al., 1998; Nautiyal et al., 2017; Srinivasan et al. 2020). Due to the process requiring crosslinking of the layers after each dip, the drying condition of the layers is important to prevent delamination for this fabrication method (Abraham et al., 1998). One of the easiest parameters to control for dip-casting process is the waiting time between the dips. Waiting time is required to allow for the

chemical reactions that crosslinks PVA molecules to occur (Chaouat, 2008). Shorter wait time between each dip prevents completion of the reactions, allowing for more partially crosslinked sites for each dip. These partial crosslinking sites can react with the newly applied layer, increasing the interlayer adhesion (Yu et al., 2013). Therefore, the drying time should be kept short for the subsequent layers to be applied before crosslinking of the previous layer is completed to increase the interlayer adhesion.

For the control, the waiting time between the dips is 15 min in the first 4 dips, 25 min in the fifth to seventh dips, and 30 min for the eighth dip and beyond (Figure 1A). Instead of increasing the dipping time throughout the fabrication process, the waiting times between dips were kept consistent for the groups comparing the effects of the interlayer adhesion. 15 W was fabricated with a flat 15-min waiting time for all waiting steps instead of the standard protocol. 30 W was made using a flat 30-min waiting time for all waiting steps. The waiting times for 15 and 30 W were kept constant so that interlayer adhesion throughout the grafts would remain consistent. There are many methods to make measurements to assess the adhesion between layers (Awaja et al., 2009). The most standard method of measuring the interlayer adhesion is the adhesive shear test (which is also known as lap shear test) (Awaja et al., 2009).

The crosslinking density and concentration of STMP are positively correlated to the stiffness of the material (Leone, 2010). As decrease in crosslinking agent can result in lowered crosslinking density, STMP concentration was decreased for one of the experimental groups (LS samples). Also, the concentration of NaOH was reported to be positively related to the crosslinking density (Lack, 2004). Therefore, 15% w/v NaOH concentration was explored to achieve decreased crosslinking density to assess the effect of low crosslinking density. However, the graft could not be fabricated using the low concentration of NaOH (Supplementary Table S1). We also speculate post-fabrication thermal drying may induce additional physical crosslinking, resulting in higher crosslinking density. To create the PVA grafts with higher crosslinking density, additional groups (60-2D, 60-4D, 18-2D, and 18-4D) were made with extra drying step post-fabrication to encourage additional physical crosslinking.

## Effects of Interlayer Adhesion on PVA Graft Compliance and Burst Pressure

As shown in Figure 3B, 15 W had significantly thinner wall than the control group. However, the wall thickness of 30 W was not significantly different from the control group. While the 15 and 30 W had consistent waiting times between the dips, control group had 15-min waiting time for the first four dips, 25-min waiting time for the 5th-7th dips, then 30-min waiting time for 8th or more dips. The data showed that the longer waiting time is needed and waiting time of the outer layer was related to wall thickness in the outer layer. When the waiting time of the outer later was decreased in 15 W, the thickness was decreased, compared to 30 W and control. However, it appears that increasing the early dip waiting time did not significantly affect the thickness.
<figure><img src="image_259.png" /></figure>

<figure><img src="image_260.png" /></figure>

<table><caption>C Longitudinal Tensile Data</caption><thead><tr><th></th><th>Ultimate Tensile Stress (MPa)</th><th>Elastic Modulus (MPa)</th></tr></thead><tbody><tr><td>C</td><td>0.494 ± 0.132</td><td>0.571 ± 0.100</td></tr><tr><td>15W</td><td>0.511 ± 0.051</td><td>0.584 ± 0.119</td></tr><tr><td>LS</td><td>0.584 ± 0.086</td><td>0.802 ± 0.456</td></tr></tbody></table>

<table><caption>D Circumferential Tensile Data</caption><thead><tr><th></th><th>Ultimate Tensile Stress (MPa)</th><th>Elastic Modulus (MPa)</th></tr></thead><tbody><tr><td>C</td><td>0.306 ± 0.124</td><td>0.690 ± 0.136</td></tr><tr><td>15W</td><td>0.313 ± 0.067</td><td>0.674 ± 0.098</td></tr><tr><td>LS</td><td>0.326 ± 0.105</td><td>1.048 ± 0.126*</td></tr></tbody></table>

<figure><img src="image_261.png" /><figcaption>**FIGURE 9** | Additional bulk mechanical properties of the selected groups. (A) Longitudinal tensile curve. (B) Circumferential tensile curve. (C) Longitudinal ultimate tensile strength and elastic modulus for each groups. (D) Circumferential ultimate tensile stress and elastic modulus. (E) Suture retention strength. *n* = 6 for tensile test results, and *n* = 9 for suture retention strength for all groups. *Indicates *p* < 0.05 with respect to the control group.</figcaption></figure>

15 W had higher adhesive shear strength than all of the other groups. While not statistically significant against the control group, 15 W showed a trend toward significantly higher adhesive shear strength than the control group (*p* = 0.057), and a significantly higher adhesive shear strength than 30 W (*p* < 0.05). This indicates that the 15 W had higher adhesion between layers than the other groups, which is in agreement with the height profiles shown in **Figures 5C–J**. As 15 and 30 W were fabricated with a flat 15 and 30-min waiting time, respectively, the significant higher shear strength in 15 W further demonstrated that the shorter waiting time would increase the interlayer adhesive strength. This could be the result of a larger quantity of partial crosslinking sites available for a shorter waiting time in 15 W. The partial crosslinking sites can then work as a bridge between the layers to allow for a larger number of interpenetrating crosslinks (Awaja et al., 2009). With longer waiting time between layers, the partial crosslinking sites can become crosslinked. This can reduce the crosslinking between the layers. Except for 15 and 30W, all of the samples had variable drying time. For the control group, LS, 60-2D, 60-4D, 18-2D, and 18-4D, the last two layers were applied with 30-min waiting intervals. This could be the reason these groups did not show statistical difference in the adhesive shear strength.

## Effects of Chemical Crosslinking Density on PVA Graft Compliance and Burst Pressure

15 W had a higher swelling ratio compared to the control group. The partial crosslinking sites could also attribute to the higher

swelling ratio. This partial crosslinking sites may have resulted in lower crosslinking density, which could account for the lower burst pressure. However, it is also important to note that the wall thickness of 15 W was significantly thinner than the control, indicating that the decrease in burst pressure could also be due to the thinner walls (Menon, 2011). When the burst pressures were normalized to the wall thickness of the samples, 15 W did not show significant difference from the control group (**Supplementary Figure S1**). This suggests that the affect from partial crosslinking on burst pressure was less significant than the influence from wall thickness.

The crosslinking of PVA was achieved by phosphoesterification, thus the content of phosphate also reflects the number of chemical crosslinking bonds formed during the fabrication. With lower concentration of STMP added in the reaction solution, it was expected that the resulting PVA grafts would display lower concentration of phosphate to indicate a lower chemical crosslinking density. However, LS showed lower phosphate content than the control group (**Supplementary Figure S1**). The lower phosphate content in conjunction with the lower swelling ratio, LS displayed higher crosslinking density compared to the control group. This is in agreement with the burst pressure normalized to wall thickness (**Supplementary Figure S1**). The burst pressure should increase with increase in crosslinking density. A potential explanation for higher crosslinking density is the presence of tighter network due to the multifunctional nature of STMP. As phosphate groups from STMP has 4 potential reactive sites, more than two reactions can occur on a single phosphate group. When less phosphate groups
are available, hydroxyl groups have less phosphate groups to react with. Phosphate groups become the limiting reagents, causing multiple phosphoesterification to occur on one phosphate. This results in tighter mesh formation, and therefore tighter crosslinking density. For the PVA crosslinked in abundance of phosphate, more phosphate could react with the hydroxyl groups. The abundance could result in phosphate population that did not react with multiple hydroxyl groups in the polymer network. These unbound phosphate groups then get washed away during the rehydration and storage, effectively resulted in a lower phosphate content in the hydrogels.

The difference in the crosslinking density among 15W, 30 W and the control, observed from the phosphate assay (Figure 5A), indicated that the waiting time in between different layers affected the chemical crosslinking density. 15 W had a significantly higher swelling ratio and significantly lower wall thickness than 30 W. Comparing the mechanical properties of 15 and 30W, 30 W had significantly lower compliance and significantly higher burst pressure. With a decrease in waiting time from 30 to 15 min between each dip, more interpenetrating crosslink occurs. The interpenetrating crosslinks contribute to the observation of increased chemical crosslinking density observed from the phosphate assay. Similar to LS, 30 W had longer waiting time between each dip, allowing more time for polymer chain to get aligned for hydrogen bonding, which may have contributed to the decreased compliance and increased burst pressure.

## Effects of Physical Crosslinking Density on PVA Graft Compliance and Burst Pressure

Although the chemical crosslinking density should not change during the additional drying due to extensive washing steps, physical crosslinking may have occurred during 60°C drying. After the STMP crosslinking of PVA, residual hydroxyl groups remain available for hydrogen bonding formation. PVA polymer chains display increased mobility at high temperature. The mobile polymer chains have a higher chance to get aligned, bringing the hydroxyl groups close to form hydrogen bonding. This process was achieved through subjecting the fabricated control groups to additional drying.

The extra-drying groups were 60-2D, 60-4D, 18-2D, and 18-4D. 60-2WD and 60-4WD are the two groups that were dried in the 60°C for 2 and 4 weeks, respectively. These two groups had thinner wall thickness than the control groups. 18-2D and 18-4D were the two groups that were dried in the 18°C for 2 and 4 weeks, respectively. The wall thickness of these groups did not differ from that of the control group. The two groups dried at 60°C had lower swelling ratio than the control group regardless of the length of drying, indicating that they may have higher crosslinking density than the control group. Drying in the 18°C did not alter the swelling ratio of the samples, suggesting that the crosslinking density of the groups dried post-fabrication at 18°C was not different from that of the control group. The compliance of the 18-2D and 18-4D were comparable to the compliance of 60-2D and 60-4D. All of the post-fabrication drying groups had lower compliance than the control group.

18-2D and 18-4D had similar burst pressure compared to the control group. However, 60-2D and 60-4D had higher burst pressure than the control group. There was no difference observed based on the length of extra drying (2 vs. 4 weeks) for both temperatures. This suggests that the additional physical crosslinking occurred during dehydration rather than the storage.

As the length of drying (2 vs. 4 weeks) did not affect the performance of the grafts, further analysis was performed using 18-2D and 60-2D only. According to the phosphate quantification assay (Figure 5A), 60-2D and 18-2D did not display any difference in total phosphate content within the polymers. FTIR did not also display any difference between the control group and the extra-drying groups (Figure 5B). DSC analysis was performed in addition to assess the degree of crystallinity and overall crosslinking (Figure 6 and Supplementary Figure S2). The DSC results suggested that 18-2D had higher crystallinity and crystallization temperature than the control group, while 60-2D showed a slightly higher fractional crystallinity than the control group. The higher crystallinity, suggesting a higher physical crosslinking, may explain the lower compliance and slight increase in the burst pressure from the post-fabrication drying step. However, while physical crosslinking density can be used to influence compliance and burst pressure, its effect might not be as significant as chemical crosslinking.

## Additional Mechanical Analysis of Selected Groups

The group that had the highest compliance (15W) and the group that had the lowest compliance (LS) were used to perform additional mechanical property analysis. Both the longitudinal and circumferential tensile data (Figures 9A–D) agreed with the compliance data in that the 15W and control group did not show a statistical difference in mechanical behavior. LS had lower compliance than the control group, therefore it was as expected that the group is less elastic than the control group (Figure 9D). Suture retention was not affected by the different fabrication conditions. To reiterate, 15W had higher interlayer adhesion but lower crosslinking density than the control group. Therefore, it could be that the increased interlayer adhesion can also increase the suture retention strength. LS also had a thinner wall thickness than the control group. LS had not significantly different interlayer adhesion and higher crosslinking density compared to the control group. This could indicate that higher crosslinking density can increase the suture retention strength as well.

## CONCLUSION

We investigated the effects of interlayer adhesion and crosslinking density on the compliance of PVA vascular grafts. We found that the higher crosslinking density resulted in higher burst pressure, but resulted in lower compliance. Also, we have found that the increase in interlayer adhesion could increase the burst pressure without sacrificing the compliance. While the findings in this
article sheds light on how to improve the performance of the PVA vascular grafts, there are still shortcomings of the PVA crosslinking method utilized here. PVA-STMP crosslinking method using dip-casting is a slow process. This limits the shapes, and therefore applications, of the grafts such fabrication method can produce. Furthermore, PVA tubular grafts are prone to kinking. Kinking could potentially affect the blood flow when the graft is implanted in the patients, resulting in negative clinical outcome. Further research should be performed to improve the fabrication process of PVA-STMP vascular grafts to increase the potential application. Also, the kink resistance of the grafts should be improved to result in better clinical outcome.

## DATA AVAILABILITY STATEMENT

The original contributions presented in the study are included in the article/Supplementary Material, further inquiries can be directed to the corresponding author.

## AUTHOR CONTRIBUTIONS

YJ performed the experiments, performed data analysis, and contributed to writing and editing of the article. EY performed data analysis, and contributed to writing and editing of the article.

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## FUNDING

The authors would like to thank the National Institute of Health (NIH RO1 HL130274-01A1), NSERC-CREATE Training in Global Biomedical Technology Research and Innovation at the University of Waterloo (CREATE-509950-2018), NSERC Canada Discovery Grant (RGPIN-2016-04043), John R. Evans Leaders Fund - Canada Foundation for Innovation (35573), and Mitacs Research Training Award for Financial Support.

## ACKNOWLEDGMENTS

The authors would like to thank Clyde Overby for editing the work.

## SUPPLEMENTARY MATERIAL

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**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2021 Jeong, Yao, Mekonnen and Yim. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
Polymeric and Composite Materials

Françoise Berzin

Prof. Françoise Berzin currently works at the Fractionnement des Agroressources et Environnement (FARE), Université de Reims Champagne-Ardenne. Françoise is a specialist of composite materials and polymer extrusion. She achieved her doctorate in 1998 at École des Mines de Paris (MINES ParisTech), followed by two postdoctorate positions at École Polytechnique de Montréal, Canada, and École des Mines de Paris (MINES ParisTech) and Arkema Partnership, respectively. From 2000 to 2017, she was an Associate Professor at the University of Reims Champagne-Ardenne, and she has been a Full Professor there since 2017.

Jayashree Bijwe

<figure><img src="image_262.png" /></figure>

Prof. Jayashree Bijwe has worked as a Professor at the Indian Institute of Technology, Delhi, India, for the past 25 years. Her current major research areas include the development of high-performance polymer composites/nanocomposites for tribological applications, such as Non-asbestos-Non-Copper- organic brake pads, dry bearings, and high-temperature adhesives, and the development of oils/nano-oils, greases for engines, gears, etc. She has published 200 research papers in international journals, 100 conference papers, and 12 book chapters. Her h Index (google scholar ) is 48, she has an i10-index of 142, and she has over 8000 citations. She has been on the editorial/advisory boards of various journals like Wear, Tribology International, STMP, Industrial Lubrication and Tribology, IOP SciNotes, etc.

Valentina Brunella

<figure><img src="image_263.png" /></figure>

Prof. Valentina Brunella received her Ph.D. in Chemistry, and, since 2002, she has been serving the Department of Chemistry at the University of Torino, first as Assistant Professor and now as Associate Professor in Industrial Chemistry. She has proven experience in the polymeric and composite material field. Her research interests are focused on the characterization of polymeric materials from different points of view: (i) structure-property relationships in polymers, copolymers, and composites; (ii) the study of different degradation processes of polymeric matrices under the effect of temperature, radiation, and so on; and (iii) chemical and mechanical characterization and modification of polymeric adhesive joints.
<figure><img src="image_264.png" /><figcaption>Valerie Desauziers</figcaption></figure>

Prof. Valérie Desauziers, Ph.D. in analytical chemistry, is a Professor at IMT Mines Ales, a French school for engineering and applied science. She is the head of the research group RIME (research for material/ environment interactions) located in Pau, France. Their objective is to study the impact of materials and associated processes (manufacturing, use, aging, recycling, etc.) on the environment, health, and sensory properties. Valerie Desauziers's research concerns the development of sampling and analytical methods for VOCs and SVOCs traces in air and at the material/air interface to study transfers of emerging pollutants and assess their impact on indoor air quality and health.

<figure><img src="image_265.png" /><figcaption>Patricia Maria Frontini</figcaption></figure>

Prof. Patricia Maria Frontini is a Full Professor at the University of Mar del Plata and a Superior Researcher of the National Research Council of Argentina. Her scientific activities concern the deformation and fracture of polymers and composites, performance under high-rate loading, yield behavior under multiaxial loading, the performance of injection molded parts, nanoindentation, and constitutive modeling. She has participated in numerous international scientific collaboration projects and technological developments for the productive sector. She has published more than 120 papers for international indexed journals and more the 200 conference papers. She has also received several scientific awards in Argentina (Houssay 2005; Ing. Santiago 2015, and Ing. Huergo 2019).

<figure><img src="image_266.png" /><figcaption>Carla Martins</figcaption></figure>
Dr. Carla I Martins is an Assistant Professor at the University of Minho and a researcher at the Institute for Polymer and Composites (IPC), Portugal. She holds a degree in Polymer Engineering (1998), a Master's in Math Applied to Mechanics (2001) from the University of Minho, and a Ph.D. in Polymer Engineering (2004) from the University of Akron, USA. She has participated in National and European industrial-oriented projects as both an investigator and principal investigator. She has supervised a number of Master's and Ph.D. students in the field of processing structure properties of polymeric materials, filled systems, and recycled/reused materials. She has (co-)authored 40 publications in national and international journals or conferences.

<figure><img src="image_267.png" /><figcaption>Véronique Michaud</figcaption></figure>

Prof. Véronique Michaud is currently an Associate Professor, Head of the Laboratory for Processing of Advanced Composites, and an Associate Dean of Engineering for Education at the Ecole Polytechnique Fédérale de Lausanne, Switzerland. She graduated in 1987 from Ecole des Mines in Paris with an engineering degree and in 1991 from MIT with a Ph.D. in Materials Engineering. She also obtained a Research Habilitation from INPG in France in 1994. After a post-doctoral research stay at MIT, she spent 3 years at Ecole Centrale in Paris to teach and research at the Laboratory for Materials, Structures, and Soils Mechanics before joining EPFL in 1997. Her fields of research address fundamental aspects of composite materials processing as well as the development of smart materials and structures, including self-healing, shape and vibration control, and tailored damping. She is also the co-founder of CompPair Technologies SA, which was created in 2020.

<figure><img src="image_268.png" /><figcaption>Hamideh Mirbaha</figcaption></figure>

Dr. Hamideh Mirbaha is currently a post-doctoral researcher at the laboratory of Bio-inspired, Bionic, Nano, Meta Materials & Mechanics at the University of Trento, Italy. She received her Ph.D. from Amirkabir University of Technology, Iran, where she studied the structure formation of PAN/CNT composite fibers during the solution spinning process. She was a visiting researcher at the University of Trento, Italy, during her Ph.D. research, characterizing the mechanical and structural properties of nanocomposite fibers produced under different shear and elongation conditions. Her main research areas also include the actuation behavior of polyacrylonitrile gel fibers, the stabilization process of PAN fibers, and conductive electrospun nanofibers.
<figure><img src="image_269.png" /><figcaption>Amulya Raj</figcaption></figure>

Dr. Amulya Raj received her Ph.D. in Material Engineering from IMT Lille-Douai, France, in 2019, prior to which she received a Master's degree in Polymer Science and Technology from VTU, India, in 2013. She is currently working as a Polymer Engineer at Eden Tech, Paris, with a focus on polymers in microfluidic applications. Her research interests include polymer chemistry, processing, and blending with a special emphasis on bio-based polymers. She has worked extensively on PLA blends for high-performance and durable applications.

<caption>Florence Saffar</caption>

Dr. Florence Saffar is a Postdoctoral Research Assistant at the University of Edinburgh in Advanced Materials for Renewable Energy Applications. She currently works on Fiber Metal Laminate (FML) development. She achieved her Ph.D. at ONERA – The French Aerospace Lab on the topic of the consolidation of thermoplastics in low-pressure situations.
<figure><img src="image_270.png" /></figure>

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# Use of Flow Modeling to Optimize the Twin-Screw Extrusion Process for the Preparation of Lignocellulosic Fiber-Based Composites

Françoise Berzin¹*, Chantal David² and Bruno Vergnes³*

¹FARE, University of Reims Champagne Ardenne, INRAE, Reims, France, ²Sciences Computers Consultants, Saint-Étienne, France, ³MINES ParisTech, PSL Research University, Sophia Antipolis, France

Thermoplastic polymers reinforced by lignocellulosic fibers are increasingly used to replace conventional composites based on carbon or glass fibers. These materials are generally prepared by dispersing the fibers into the polymer matrix by melt mixing in a twin-screw extrusion process. However, during the process, a significant breakage occurs, leading to a reduction in the length, and diameter of the bundles and/or individual fibers. As the mechanical properties of the composite depend, among other, on the fiber morphology, it is important to understand and control the breakage mechanisms during the compounding process. In this paper, we show how the use of a thermomechanical model of twin-screw extrusion coupled with evolution laws of the fiber dimensions makes it possible to calculate the variation in the length and the diameter of the fibers, according to the local values of the flow parameters (shear rate, residence time, temperature, etc.). It is thus possible to define the best processing conditions (screw speed, feed rate, and barrel temperature) or the best screw profile to limit fiber degradation and prepare composites with optimal properties.

**Keywords:** lignocellulosic fibers, extrusion, modeling, breakage, optimization

## INTRODUCTION

Due to their interesting properties, such as lightness, low cost, biobased origin, and satisfactory mechanical properties, lignocellulosic fibers are increasingly used for the preparation of thermoplastic matrix composites (Mohanty and Drzal, 2005; Hodzic and Shanks, 2014; Campilho, 2015). For short fibers (a few mm long) and thermoplastic matrices, the preferred process is melt mixing, i.e., compounding followed by extrusion or injection molding. In a first step, the fibers are mixed with the polymer matrix to obtain composite pellets, generally using a twin-screw extruder. In a second step, these pellets are processed by extrusion or injection molding to obtain the desired final shape. During the compounding step, many problems arise: first, when subjected to high stresses and strains, the fibers break. Not only the length, but also the diameter, of the initial bundles or the individual fibers can be considerably reduced. Second, the dispersion of the fibers into the matrix may be heterogeneous, with the existence of agglomerates. Third, if the temperature rises above about 200°C, thermal degradation can occur (Tajvidi and Takemura, 2010). Therefore, it is clear that the control of the extrusion conditions (screw profile, screw rotation speed, feed rate, etc.) is essential to prepare composites with optimal properties.

**OPEN ACCESS**

**Edited by:**

Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**

Jose Antonio Covas,
University of Minho, Portugal
Margaret J. Sobkowicz,
University of Massachusetts Lowell,
United States

**\*Correspondence:**

Françoise Berzin
francoise.berzin@univ-reims.fr
Bruno Vergnes
bruno.vergnes@mines-paristech.fr

**Specialty section:**

This article was submitted to
Polymeric and Composite Materials,
a section of the journal
Frontiers in Materials

**Received:** 15 May 2020

**Accepted:** 15 June 2020

**Published:** 24 July 2020

**Citation:**

Berzin F, David C and Vergnes B (2020) Use of Flow Modeling to Optimize the Twin-Screw Extrusion Process for the Preparation of Lignocellulosic Fiber-Based Composites. Front. Mater. 7:218. doi: 10.3389/fmats.2020.00218
The role of the processing conditions on the fiber state and on the final properties of composites has been studied by many authors (Iannace et al., 2001; Le Baillif and Oksman, 2009; Quijano-Solis et al., 2009; Barkoula et al., 2010; Mano et al., 2010; Beaugrand and Berzin, 2012; El Sabbagh et al., 2014; Gonzalez-Sanchez and Gonzalez-Quesada, 2015; Feldmann et al., 2016). In recent years, we have developed a systematic study of the breakage of lignocellulosic fibers under thermomechanical stresses, based on rheo-optical observations (Le Duc et al., 2011; Castellani et al., 2016), and experiments carried out in internal mixers (Di Giuseppe et al., 2017) and twin-screw extruders (Berzin et al., 2017, 2018, 2019, 2020). These studies were conducted with different thermoplastic matrices (polypropylene and polybutylene succinate) and various lignocellulosic fibers (flax, hemp, sisal, miscanthus, and pineapple leaf fibers). The main result is that fiber breakage (i.e., reduction in the length and diameter of bundles and individual fibers) occurs primarily by fatigue, after numerous bends, unlike glass fibers which, due to their rigidity, break when a critical stress is reached (Durin et al., 2013). Consequently, the parameter which controls the kinetics of breakage is the strain accumulated by the fiber during the flow along the twin-screw extruder. This mechanism depends obviously on the properties of the fiber (plant origin, composition, water content, etc.) but also on its dimensions (longer fibers break more easily than short ones) and on the properties of the matrix (viscosity, affinity with fiber, etc.).

As the reinforcing properties of the fibers in the composite are among other controlled by their length (or their aspect ratio), it is necessary to limit breakage during compounding by twin-screw extrusion. However, it remains long and difficult by trial and error to find the optimal conditions, providing the maximum output while preserving the fiber length. In this paper, we will show how the use of flow modeling may allow to solve this delicate problem and to rapidly test many different solutions to select the most interesting.

# MATERIALS AND METHODS

## Composite Preparation

As noted in the introduction, we focus on the compounding step, where the fibers are dispersed into the matrix by using a co-rotating twin-screw extruder. The laboratory scale extruder used by Berzin et al. (2017) (Cletral BC21, Firminy, France) was selected as an example of application. It has a diameter of 25 mm and a length of 900 mm. The screw profile is shown in Figure 1. It comprises nine barrel elements (numbered from 1 to 9, from the hopper to the screw end) and, in addition to the screw conveying elements, it includes a left-handed element in zone 3 to melt the polymer matrix and two mixing blocks in zones 6 and 8, to disperse the fibers. The first one has five kneading discs, staggered by 90°, while the second one is more restrictive, with five kneading discs staggered by -45°. The matrix is introduced in zone 1, melted in the left-handed element (zone 3), then the fibers are added in zone 4.

The matrix is that used in the paper by Berzin et al. (2017). It is a mix of polypropylene (PP) and maleic anhydride grafted

polypropylene (PP-g-MA) used as compatibilizer. Indeed, PP-g-MA is commonly used to improve the affinity between the hydrophilic fibers and the hydrophobic matrix (Feng et al., 2001; Bos et al., 2006). The PP-g-MA/fiber ratio is 1/10 by weight, as recommended in the literature (Le Moigne et al., 2011). Among the various fibers studied by Berzin et al. (2017), we have selected flax fibers for this paper. They were provided by FRD® company (Fibres Recherche Développement, Troyes, France) and have the following average dimensions: length L = 4 mm, diameter D = 242 µm, aspect ratio L/D = 18.8. They contain 86.5 wt% of cellulose, 8.6 wt% of hemicellulose and 4.9 wt% of lignin.

## Breakage Mechanisms

As indicated in the Introduction, the breakage mechanisms of flax fibers have been characterized and quantified in previous studies (Berzin et al., 2017; Di Giuseppe et al., 2017). For a fixed composition (type of fiber and matrix), it has been shown that the length L and the diameter D vary with the cumulated strain Γ, according to exponential functions. Local strain (without dimension) is the product of the local shear rate by the corresponding local residence time. It is then cumulated from the introduction of the fibers into the extruder up to the sampling point considered. The following expressions have been proposed and validated for various types of fibers and matrices (Berzin et al., 2017, 2018, 2019, 2020):

$$L = L_{\infty} + (L_{0} - L_{\infty}) \exp(-k_{L} \Gamma) \quad (1)$$

$$D = D_{\infty} + (D_{0} - D_{\infty}) \exp(-k_{D} \Gamma) \quad (2)$$

where $L_0$ (respectively, $D_0$) is the initial fiber length (respectively, diameter), $L_\infty$ (respectively, $D_\infty$) is its ultimate length (respectively, diameter), $\Gamma$ is the cumulated strain and $k_L$ (respectively, $k_D$) is a constant, quantifying the breakage kinetics. These constants depend on the properties of the fiber and the matrix. In the case of selected 4 mm flax fibers, the following values were obtained by Berzin et al. (2017): $L_0 = 4$ mm, $L_\infty = 1.2$ mm and $k_L = 1.1 \times 10^{-3}$, and $D_0 = 242$ µm, $D_\infty = 100$ µm, and $k_D = 2.8 \times 10^{-3}$. The corresponding experimental data and the theoretical curves provided by Equations (1) and (2) are illustrated in Figure 2.

## Flow Modeling

The strain cannot be measured experimentally on a twin-screw extruder. Therefore, it is mandatory to use a model to calculate it and to assess changes in length and diameter of the fibers during the process. In this paper, the Ludovic® software developed by Vergnes et al. (1998) and marketed by SCC company (Saint Etienne, France) was used. Ludovic® is a model based on continuum mechanics, which solves the Stokes and energy equations within the framework of simplified geometry and kinematics. It simulates the whole extrusion process, from the introduction of solid polymer pellets in the hopper to the die exit, and calculates the variations along the screws of the main flow parameters, such as pressure, temperature, residence time, shear rate, stress, strain, etc. The simulation is based on a one-dimensional analysis. The screws are divided into elementary
<figure><img src="image_271.png" /><figcaption>FIGURE 1 | Screw profile. The restrictive elements are in gray.</figcaption></figure>

<figure><img src="image_272.png" /></figure>

<figure><img src="image_273.png" /><figcaption>FIGURE 2 | Changes in length (A) and diameter (B) of 4 mm flax fibers as a function of cumulative strain. Symbols are experimental points, the lines are fits using (Equations 1, 2). Reprinted from Berzin et al. (2017) with permission.</figcaption></figure>

sub-elements in which mass and thermal balances are written. In each type of element (left- and right-handed screw elements, blocks of kneading discs, etc.), average values of flow parameters are computed. The computation starts at the die exit and, element by element, temperature, and pressure are computed

by an iterative backward procedure. We refer to the previous publication by Vergnes et al. (1998) for more details. For 20 years, the Ludovic® software has been extensively validated, experimentally (Carneiro et al., 2000) and by comparisons with more elaborate 3D models (Durin et al., 2014).

## Optimization Procedure

The objective is to define the best processing conditions to prepare composites with the screw profile presented in Figure 1, based on the products defined in section Composite Preparation, i.e., a PP/PP-g-MA matrix reinforced by 20 wt% of flax fibers. The best conditions mean reaching the maximum flow rate while keeping the fibers as long as a possible, i.e., at the lowest cumulated strain. However, other constraints must be taken into account.

As explained in the Introduction, the fibers must be correctly distributed into the matrix. There are many indices to characterize the distributive mixing, but the simplest is to consider the variance $\sigma^2$ of the residence time distribution (RTD) (Manas-Zloczower, 2009). The RTD is usually represented by a function $E(t)$ which gives the extruder response to an input pulse (Danckwerts, 1953). By definition, $E(t)dt$ is the fraction of material leaving the extruder with a residence time between $t$ and $t + dt$. It is such as:

$$ \int_{0}^{\infty} E(t)dt = 1 \qquad (3) $$

The mean residence time $\bar{t}$ is the first moment of the distribution:

$$ \bar{t} = \int_{0}^{\infty} tE(t)dt \qquad (4) $$

The variance $\sigma^2$ is the second moment centered on the mean:

$$ \sigma^2 = \int_{0}^{\infty} (t - \bar{t})^2 E(t)dt \qquad (5) $$

$\sigma^2$ represents the width of the RTD. In extrusion, it is directly linked to the intensity of the back-mixing. As an example for the application illustrated below, an arbitrary value has been chosen, i.e., a variance >200 s² will be requested.

To avoid thermal degradation, the maximum temperature encountered during the process should not exceed 200°C (Gassan
and Bledzki, 2001). As for the strain, this maximum value cannot be measured as it can be reached somewhere inside the extruder, typically in a block of kneading disc. Again, modeling is necessary to obtain this information.

Finally, the last constraint is the maximum loading that can be fed to the extruder. In practice, the fiber feeding is a real challenge because of their very low apparent density, close to 0.1, leading to a large volume to introduce. The maximum feeding capacity can be estimated as follows. The free cross section of the extruder S in the feeding zone can be written:

$$S = S_B - 2S_S \quad (6)$$

where $S_B$ is the cross section of the eight-shaped barrel and $S_S$ that of the screw. $S_B$ and $S_S$ can be easily calculated using the expressions proposed by Booy (1978):

$$S_B = 2(\pi - \psi)R_e^2 + C_l R_e \sin\psi \quad (7)$$

$$S_S = n[\psi C_1^2 - C_l R_e \sin\psi] + \frac{n\alpha}{2} [R_e^2 + (C_l - R_e)^2] \quad (8)$$

where $R_e$ is the screw radius, $C_l$ is the centerline distance, n is the number of flights, and $\psi$ and $\alpha$ are two angles, respectively, defined as:

$$\cos\psi = \frac{C_l}{2R_e} \quad (9)$$

$$\alpha = \frac{\pi}{n} - 2\psi \quad (10)$$

The free cross section S moves downstream at a velocity V equal to:

$$V = \frac{NB}{60} \quad (11)$$

where N is the screw speed (expressed in rpm) and B the screw pitch.

The theoretical free volume by time unit is then $Vol = S^{NB/60}$. However, as the fibers fall out of the hopper, not all of this volume is directly accessible. Only the upper part of the cross section of the channel can be reached by the fibers. Moreover, part of the channel is already filled by the molten polymer. We can thus estimate that the real free volume is only about a third of the theoretical one. This approximation is sufficient because we are simply looking for an order of magnitude, not an exact value. It was checked that this order of magnitude agreed with the experimental observations. The maximum mass flow rate of the fibers will then be approximated as:

$$Q_f \max \approx \rho_a S \frac{NB}{180} \quad (12)$$

where $\rho_a$ is the apparent density of the fibers, typically close to 0.1.

<figure><img src="image_274.png" /><figcaption>FIGURE 3 | Changes in maximum temperature with screw speed for different feed rates.</figcaption></figure>

To summarize, the following constraints have to be respected for the optimization:

- maximum flow rate;

- minimum cumulated strain $\Gamma$;

- variance of RTD higher than 200 s²;

- maximum temperature <200°C;

- fiber flow rate lower than $Q_f \max$.

## RESULTS AND DISCUSSION

The optimization we propose is not a formal optimization, using specific algorithms, as we have done in other publications (Gaspar-Cunha et al., 2005, 2011), but a practical one, allowing to quickly define an optimal processing window. It will be conducted in two steps. First, the best processing conditions will be defined for the screw in Figure 1, according to the constraints defined in section Optimization Procedure. Second, we will show how changing the screw profile can help to improve the first results.

### Optimization of Processing Conditions

To define the best processing conditions, simulations with Ludovic® software are performed by varying the total flow rate (PP/PP-g-MA + 20 wt% flax fibers) from 2 to 6 kg/h and the screw speed from 50 to 200 rpm. For each simulation, the values of the cumulated strain $\Gamma$, the RTD variance $\sigma^2$ and the maximum temperature $T_{max}$ are recorded. The results are presented in Figures 3–5.

The maximum temperature $T_{max}$ increases a lot with the screw speed, due to viscous dissipation (Figure 3). It can reach values of 30–40°C above the barrel temperature (180°C), capable of degrading not only the fibers but also the matrix. At a fixed screw speed, $T_{max}$ decreases slowly with the feed rate. In order to
<figure><img src="image_275.png" /><figcaption>**FIGURE 4** | Changes in RTD variance with screw speed for different feed rates.</figcaption></figure>

<figure><img src="image_276.png" /><figcaption>**FIGURE 5** | Changes in cumulated strain with screw speed for different feed rates.</figcaption></figure>

keep the temperature below the maximum value of 200°C, low screw speeds (in the range 110–150 rpm) are necessary.

The variance of the RTD $\sigma^2$, selected as the distributive mixing index, is shown in **Figure 4**. It decreases slightly with the screw speed, but is mainly affected by the flow rate. These effects of screw speed and feed rate on the RTD are well-known in twinscrew extrusion (Poulesquen and Vergnes, 2003): an increase in screw speed shifts the RTD toward shorter times, without affecting its shape, while an increase in feed rate also shifts the RTD toward shorter times, but mainly narrows its width. To comply with the constraint of RTD variance >200 $\sigma^2$, low flow rates are mandatory: the limit appears between 3 and 4 kg/h.

<figure><img src="image_277.png" /><figcaption>**FIGURE 6** | Processing window on a screw speed/feed rate map. Hatched regions are forbidden by the various criteria.</figcaption></figure>

<figure><img src="image_278.png" /><figcaption>**FIGURE 7** | Calculated changes in fiber length and diameter along the screw profile in the optimized conditions.</figcaption></figure>

Finally, the values of the cumulated strain $\Gamma$ are presented in **Figure 5**. $\Gamma$ increases linearly with the screw speed and decreases with the feed rate. Indeed, it can be shown that $\Gamma$ is proportional to the ratio $N/Q$ (Lafleur and Vergnes, 2014). Therefore, to limit the breakage of the fibers, it is necessary to favor low speeds and high feed rates.

If we consider these different results, it is clear that certain conditions appear to be contradictory. To have a clearer view of the situation, plots of isovalues of the limiting parameters $T_{max}$, $\sigma^2$, and $Q_{fmax}$ are plotted on a flow rate/screw speed map (**Figure 6**). This shows that the processing windows is quite narrow: the screw speed is limited at 100–110 rpm by the maximum temperature and the flow rate below 4 kg/h by the RTD variance. Moreover, another limit is added by the
<figure><img src="image_279.png" /><figcaption>**FIGURE 8** | Modified screw profile. The restrictive elements are in gray.</figcaption></figure>

maximum feeding rate of the fibers. In conclusion, it appears that the optimal conditions (circle in **Figure 6**) are $Q = 3.6$ kg/h at $N = 95$ rpm. These values provide a maximum temperature of $194^{\circ}$C, reached in the second block of kneading discs, a RTD variance of $200$ s$^2$, and a cumulated strain of 2,157. For these processing conditions, it is now possible to use (Equations 1, 2) to calculate the variations of the fiber dimensions along the screws. For example, **Figure 7** shows the change in fiber length and diameter under the optimized conditions. Fibers of 1.46 mm in length and $100 \ \mu m$ in diameter are obtained at the die exit.

## Optimization of Screw Geometry

It can be seen in **Figure 7** that the fibers undergo a significant reduction in both the length and the diameter, even though the processing conditions have been optimized. Degradation occurs rapidly, along the first screw conveying elements, as soon as the fibers are introduced into the extruder. Therefore, to limit this problem and reduce the fiber attrition, the simplest way is to move the fiber introduction and the two blocks of kneading discs downstream, as shown in **Figure 8**. However, a certain distance along the screws remains necessary to ensure a good distribution of the fibers into the matrix. By repeating the same procedure as in the previous paragraph (different flow simulations under various conditions of screw speed and feed rate) and with the same constraints (on $T_{max}$, $\sigma^2$, and $Q_{fmax}$), the processing window for this new configuration can be established. The optimal conditions are very close to those of the first profile, with a flow rate of 3.6 kg/h and a screw speed of 100 rpm. However, as the fibers are introduced later, the cumulated strain is lower, here 1,820, and, as expected, the fibers are a little longer: 1.58 mm (i.e., a gain of 8%). Nevertheless, the improvement is not really significant and further modifications to the geometry would be necessary for better results. As the main limitation of the flow rate is the RTD variance, we can imagine to increase this

parameter by increasing the length of the first block of kneading discs, staggered at $90^{\circ}$ (50 mm instead of 25 mm). In doing so, the numerical simulation shows that at 120 rpm and always with the same constraints, the feed rate can be increased up to 4.8 kg/h, i.e., a gain of 33%. However, the cumulated strain remains around 2,000, indicating no further improvement in the fiber length. Based on the previous results of Berzin et al. (2017), starting from initial longer fibers could be a solution: indeed, for flax fibers 12 mm long, the same processing conditions (4.8 kg/h, 120 rpm) would lead to final fibers about 3 mm in length.

## CONCLUSION

In this paper, it has been shown that the use of a model of the flow along a twin-screw extruder coupled with evolution laws for the dimensions of lignocellulosic fibers can be a very useful tool for optimizing the preparation of natural fiber composites. In a very short time (each calculation requires only a few seconds on a PC), it allows one to test many possible conditions and, according to certain constraints, to optimize the processing conditions and the screw profile. Similarly, flow modeling could also be a privileged tool for solving the often difficult problem of scale-up, from laboratory to industrial scale.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

## AUTHOR CONTRIBUTIONS

FB, CD, and BV contributed equally to the paper, by performing the simulations, discussing the results, and writing the text. All authors contributed to the article and approved the submitted version.

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**Conflict of Interest:** CD was employed by the company SCC.

The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Berzin, David and Vergnes. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_280.png" /></figure>

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# Roles of Size, Shape, Amount, and Functionalization of Nanoparticles of Titania in Controlling the Tribo-Performance of UHMWPE Composites

Meghashree Padhan, Gourab Paul and Jayashree Bijwe*

Centre for Automotive Research and Tribology (Formerly ITMMEC), Indian Institute of Technology Delhi, Hauz Khas, India

The present research is based on development and performance evaluation of ultra-high molecular weight polyethylene composites with Titania particles in various amounts and of two sizes, including nano-size. It also included functionalization of nanoparticles (NPs) with a siloxane group to enhance adhesion with the matrix, which would help to reduce wear. Thus, the work had several objectives such as to investigate the effect of size (nano vs. micro) and determine the amount that delivers the best performance from among a series of composites with different Titania particle amounts [0–8% in the case of microcomposites (MCs) and 0–4% in the case of nanocomposites (NCs)]. It was also of interest to examine the effect of siloxane-functionalization of NPs [since they performed significantly better than microparticles (MPs)] on the performance of NCs. It was observed that for adhesive and erosive wear modes, both types of particles (MPs and NPs) proved detrimental in spite of the enhanced hardness of the composites, essentially because of the irregular shape of particles, which altered the wear mechanisms adversely. Nanoparticles proved less detrimental, while F-NPs (functionalized NPs) proved marginally detrimental. In abrasive wear mode, both proved beneficial, NPs being more effective than MPs, and F-NPs proved the best. The studies revealed how the shape of particles influences the performance of composites in various wear modes.

**Keywords:** UHMWPE, TiO₂, siloxane treatment, abrasive and adhesive wear, erosion

## INTRODUCTION

Ultra-high molecular weight polyethylene (UHMWPE) and its composites are being widely used for applications like producing bottles for runners, chain guides, and lining for coal chutes and also as bearing material in agricultural and mining equipment due to their excellent wear resistance coupled with their low coefficient of friction (μ), along with ease of availability and cost viability. Moreover, UHMWPE is extensively used in a composite form as a bearing material for implants (joint replacements for the hip, knee, and ankle) (Wang et al., 2007; Ge et al., 2009; Holland et al., 2018; Gao et al., 2019; Ortiz-Hernández et al., 2019). It shows excellent wear resistance (W~R~) in almost all wear modes (Lee, 1985) along with ultra-low μ, high specific strength and modulus,

**OPEN ACCESS**

**Edited by:**

Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**

Sergey Panin,
Institute of Strength Physics
and Materials Science (ISPMS SB
RAS), Russia

Abdul Samad Mohammed,
King Fahd University of Petroleum
and Minerals, Saudi Arabia

Z. Peng,
University of New South Wales,
Australia

**\*Correspondence:**

Jayashree Bijwe
jbijwe@gmail.com

**Specialty section:**

This article was submitted to
Polymeric and Composite Materials,
a section of the journal
Frontiers in Materials

**Received:** 27 March 2020

**Accepted:** 04 June 2020

**Published:** 28 July 2020

**Citation:**

Padhan M, Paul G and Bijwe J (2020) Roles of Size, Shape, Amount, and Functionalization of Nanoparticles of Titania in Controlling the Tribo-Performance of UHMWPE Composites. Front. Mater. 7:205. doi: 10.3389/fmats.2020.00205
toughness, very good dampening capability, excellent resistance to chemicals, etc. Efforts are being made to maximize its W<sub>R</sub>, especially by selecting biocompatible fillers like titania, zirconia, etc., and their use as nanoparticles (NPs). The final performance of nanocomposites (NCs) depends on the type of matrix and NPs, their amount, size, shape, filler-matrix adhesion, etc. A lot has been reported on the exploration of various types and amounts of hard NPs (non-bio-compatible) to enhance the W<sub>R</sub> of UHMWPE (Chang et al., 2013; Xu and Tangpong, 2013; Chukov et al., 2014; Sharma et al., 2015, 2016; Kumar et al., 2017).

Xu and Tangpong (2013) presented the current trends of research in the tribology of polyethylene-based NCs along with the type of nano-fillers used and the processing approaches. It was stated that agglomeration takes place in a higher concentration of NPs, resulting in deterioration of wear performance.

Few papers have reported on the surface treatment/ functionalization of fillers/fibers to improve the performance with UHMWPE. Chukov et al. (2014) studied the effect of nitric acid treatment of carbon fibers leading to enhanced adhesion with the matrix and hence improved mechanical properties. Tribo-properties, however, were not studied. Chang et al. (2011) reported improvement in mechanical and antibacterial properties of UHMWPE due to reinforcement with silane-treated zinc oxide particles (<1 µm) compared to untreated ones. However, no tribo-studies were reported. Kandeva et al. (2018) reported the effect of multi-walled carbon nano-tubes (CNTs) untreated and modified with electro-less nickel plating. Tribo-studies were done in a thumb disk tribo-setup in abrasive wear mode. Modified CNTs showed higher W<sub>R</sub> as compared to the unmodified ones.

A lot of patents are available on the applications of UHMWPE in composite form for the biomedical area (Senatov et al., 2014; Coker, 2019; Gopanna et al., 2019; Salari et al., 2019; Saravanan and Devaraju, 2019; Duan et al., 2020)<fnref n="1" />,<fnref n="2" />. Several researchers have recommended their composites for such applications. Duan et al. (2020) studied the thermal stability and mechanical properties of irradiation-cross-linked graphene oxide/UHMWPE composites with vitamin E. The crystallinity, thermal stability, hardness and scratch resistance, and oxidation resistance of the composites improved as compared to UHMWPE, but at the cost of a slight deterioration in mechanical properties. The fatigue strength, compressive and tensile properties, and wear resistance of the modified composites were not investigated. These composites were claimed to have bright application prospects in the field of artificial joint replacement owing to some further investigations. Senatov et al. (2020) prepared a novel biomimetic highly porous UHMWPE/Hydroxyapatite (HA) (40 wt.%) scaffold for reconstructive surgery with a porosity of 85 vol.%. This biocompatible highly porous UHMWPE/HA scaffold, used separately or in combination with proteins for the reconstruction of non-load bearing parts of bones, was reported to solve the problems associated with differences between implant

architecture and trabecular bone, low osteointegration, and bio-
inertness. Saravanan and Devaraju (2019) tribo-investigated an γ-
irradiated composite of UHMWPE (80 wt.%), alumina (3 wt.%),
HA (5 wt.%), and Chitosan (12 wt.%) and reported a 1.3 times
increase in tribo-performance due to irradiation. The composites
were targeted at biomedical applications. Salari et al. (2019)
studied the effect of nano-zirconia (2, 6, and 10 wt.%) in
HP (10 wt.%)-reinforced UHMWPE composites and reported
significantly improved mechanical properties and wear resistance
in comparison with HAp alone and furthermore eliminated the
concern of wear debris and adverse biological responses caused
by osteolysis. Zirconia also demonstrated exceptional biological
responses in the presence of MG-63 osteoblast cell cultures,
showing no cytotoxicity and high alkaline phosphatase activity
along with fully spread osteo-cells. Hence, the authors claimed
it as a promising material for overcoming the shortcomings
of the conventionally used UHMWPE as an implant liner.
Gopanna et al. (2019), in their overview of polyethylene and
polypropylene composites for biomedical applications, suggested
that UHMWPE composites reinforced with NPs, nanotubes, and
nanofibers should be used in orthopedic bearing applications.
Chemically cross-linked UHMWPE filled with particles of Al, Cu
Fe, and reinforced with quartz particles of micron size have been
developed for orthopedic applications.

Titanium dioxide (TiO<sub>2</sub>)/Titania is a hard filler with excellent resistance to UV and chemicals and is also known for bio-compatibility (Hashimoto et al., 2003; Gutwein and Webster, 2004; Lei et al., 2010; Sabetrsekh et al., 2011; You et al., 2014; Celebi Efe et al., 2019). **Supplementary Appendix A** gives details on various available papers on UHMWPE-TiO<sub>2</sub> composites with special reference to NPs and their functionalization (Hu et al., 2016; Celebi Efe et al., 2017; Yu et al., 2017). It seems that only two papers dealt with NPs, and these did not look into tribo-effects (Ortiz-Hernández et al., 2019; Duan et al., 2020). Some papers (Duan et al., 2020; see footnote 1) reported on the functionalization of microparticles (MPs) but not in detail by reporting tribology. Thus, it was concluded that in-depth studies of UHMWPE-Titania composites (MCs and NCs) are required with the following aims, since they are not met in the literature.

* Assess the potential of NPs vs. MPs to improve the tribo-performance in various wear modes.

* Determine the influence of increasing amounts of NPs (realistic amount) and MPs as a means to find the optimum amount for the best W<sub>R</sub> value.

* Determine the effect of functionalization of NPs with siloxane treatment on tribo-performance in various wear modes.

With these aims in view, three series of UHMWPE-Titania composites – the first with MPs (0–8%), the second with NPs (0–4%), and the third with functionalized NPs (0–4%) – were developed and characterized for tribo-performance in three wear modes, viz., adhesive, abrasive, and erosive, and the findings are presented in subsequent sections.

<footnote n="1">¹https://ryortho.com/breaking/b-one-total-hip-cleared-by-fda/</footnote>

<footnote n="2">²https://wwwmaterialsforengineering.co.uk/engineering-materials-news/researchers-create-radiopaque-polyethylene-for-the-visualisation-of-medical-implants/110699/</footnote>
# MATERIALS AND METHODS

## Materials and Characterization of Particles

Ultra-high molecular weight polyethylene was procured in a particulate form (≈40 µm) from Bhilwara Polymers Ltd., Rajasthan; details of its properties are given in Table 1.

Titania (TiO₂) nano (10–25 nm)- and MPs (30–40 µm) were procured from Nano Shell LLC United States. The size was verified with SEM (scanning electron micrograph; Zeiss EVO MA 10) and TEM (transmission electron micrograph; JEOL JEM-1400 Plus), as shown in **Figures 1a,b**, respectively. It seems that MPs were of highly irregular shape and NPs were of 40–50 nm size and also of irregular shape. **Figure 1-II** shows the SEM and EDAX micrographs of selected specimens (V₈M, V₂M, V₂N, and F₂N) cut to expose the inside transverse section. This was done to determine whether particles are uniformly distributed in the bulk of the composites. Dot maps of Ti represent the distribution of TiO₂ particles irrespective of size. For F₂N, additional dot mapping was done for Si element to ensure the presence of siloxane treatment on the particles.

The XRD-X-ray diffraction spectrum (Phillips XPERT-PRO) of NPs was collected to verify the size, and the data are shown in **Figure 1-III**.

The crystallite size of NPs was determined using the Debye-Scherrer method using the following equation (Holzwarth and Gibson, 2011).

$$t = \frac{0.9\lambda}{B \cos \theta} \qquad (1)$$

where λ is the characteristic wavelength, B is the full width at half maximum, and θ is the midpoint of the peak. The XRD shows coherent diffraction regions that imply the size of the crystallites for the respective diffraction peak. A particle may be made up of several different crystallites or just one crystallite. Sometimes, it also corresponds to the size of the grain s of an analyte sample, or the thickness of a polycrystalline thin film or bulk material. **Supplementary Appendix A** collects the data from XRD.

Overall, it was confirmed that the approximate crystallite size of the particles was in the range of 23–27 nm as observed from the first three peaks.

## Experimental Work
### Surface Modification of Nanoparticles

For siloxane treatment of NPs, Amino propyl tri-ethoxy silane (APTES) was used (procured from Alpha Aesar). The basic aim of the surface modification of TiO₂ NPs was to impart siloxane

functional groups to the surface of particles, which would bridge with matrix, resulting in enhanced filler-matrix adhesion in the composites. It was anticipated that this might lead to improved tribological properties (less wear). Surface modification of TiO₂ with APTES was done as per a reported process as shown in **Figure 2** (Chukov et al., 2014) using reflux reaction and mechanical stirring, with xylene as a medium and at 50°C in a silicon oil bath for varying durations, as shown by **Table 2**.

The particles were washed using xylene, followed by ethanol two times, and finally with acetone. They were then filtered using a REMI PR-24 centrifuge followed by drying under vacuum. Fourier Transform Infrared (FTIR) Spectroscopy was done in attenuated transmission reflectance (ATR) mode in a Nicolet IS 50 for confirmation of the presence of siloxane treatment on the particles as well as for optimization of the time duration of the treatment. The analysis was performed for an iteration of 32 scans for accuracy of results (Table 2).

## Development of UHMWPE Composites

Uniform dispersion of micro- and nano-TiO₂ particles is extremely important. De-agglomeration of NPs was performed using a Crome Tech Probe Sonicator (UP-800 ultrasonic processor). An accurately weighed amount of NPs was added to double-distilled water in installments, followed by sonication for two minutes (pulse time 2 s, and pulse off for 5 s, at 800 W) after each installment, the total time of sonication being 30 min. A weighed amount of UHMWPE was then added to this suspension in installments followed by sonication to achieve homogeneous dispersion of NPs and UHMWPE in double-distilled water. A similar procedure was employed in the case of microcomposites (MCs) to maintain uniformity.

Nano- and microcomposites of UHMWPE with TiO₂ were developed by using a compression molding technique under temperature (180°C) and pressure (12–15 bar). The mold was fast cooled using water cooling. The compositions of the composites developed are shown in Table 3.

## Characterization of Composites
### Density

The density of the composites was measured to characterize them, since density is the first basic property. It also helped to gain an overall view of the void content in the composites and also to calculate the specific wear rate.

(a) Density was measured by the standard Archimedes principle following ASTM D792 using a Mettler Toledo weighing balance as per Eq (2). Acetone was used as the liquid, having a density of 0.79 g/cm³ at 20°C.

$$\rho = (\rho_{\text{liquid}} \times W_{\text{air}}) / (W_{\text{air}} - W_{\text{liquid}}) \qquad (2)$$

where ρ is the density of the composite, ρliquid is the density of the liquid used, and Wair and Wliquid are the weights of the composite in air and liquid.

The rule of mixtures was used to calculate the theoretical density of the composites using the following equation:

$$\rho_t = 1 / (W_f / \rho_f + W_m / \rho_m) \qquad (3)$$

<table><caption>**TABLE 1** | Properties of UHMWPE as per the supplier's data.</caption><thead><tr><td>Property</td><td>Value</td></tr></thead><tbody><tr><td>Density (g/cm³)</td><td>0.93</td></tr><tr><td>Tensile strength at yield (MPa)</td><td>21</td></tr><tr><td>Tensile strength at break (MPa)</td><td>48</td></tr><tr><td>Hardness, Shore D</td><td>62–66</td></tr><tr><td>Crystalline melting range, powder, °C</td><td>138–142</td></tr></tbody></table>
<figure><img src="image_281.png" /><figcaption>FIGURE 1 | Continued</figcaption></figure>
<figure><img src="image_282.png" /><figcaption>**FIGURE 1** | Micrographs of TiO₂ particles; (1a) SEM of microparticles and (1b) TEM of nanoparticles; 1-II – SEM and EDAX micrographs of representative composites showing uniform distribution of filler particles; 1-III – X-ray diffraction pattern for nano-TiO₂ particles.</figcaption></figure>

<figure><img src="image_283.png" /><figcaption>**FIGURE 2** | Schematic of treatment of TiO₂ NPs with APTES.</figcaption></figure>

<table><caption>**TABLE 2** | Treatment of NPs of TiO₂ (1.5 g) with APTES (1.5 ml) in xylene as a medium (50 ml) at 50°C for various durations and the chemical changes confirmed from FTIR studies.</caption><thead><tr><td>Treatment</td><td>Heating time (h)</td><td>Observation</td></tr></thead><tbody><tr><td>Scheme 1</td><td>10</td><td>No change</td></tr><tr><td>Scheme 2</td><td>48</td><td>Low-intensity peak of -Si-O-; -O-Ti-O- peak</td></tr><tr><td>Scheme 3</td><td>108</td><td>Sharp and intense -Si-O- peak; -Ti-O-Si- peak</td></tr></tbody></table>

where $ρ_m$ and $ρ_f$ are the theoretical density of the matrix (UHMWPE; 0.93 g/cc) and the filler (Ti; 3.9 g/cc), respectively. $W_m$ and $W_f$ are the respective weight fractions.

(b) Void content.

The void content of the composites was calculated from the following equation:

$$ \text{Void content (\%)} = \frac{\rho_{th} - \rho_{c}}{\rho_{th}} \times 100 \quad (4) $$

where $ρ_{th}$ and $ρ_c$ are the calculated and theoretical densities of the composite as obtained using Eqs 2 and 3.

## Vickers Micro-Hardness

TiO₂, being a ceramic material, has an entirely different set of properties, especially hardness, as compared to the softer polymeric phase. It was necessary to examine the extent of its implications for the hardness properties of UHMWPE since hardness is generally well-correlated with abrasive wear performance. The Vickers hardness of the composites was measured as per ASTM E92-82. The composites were polished against 1000 grade SiC abrasive paper. A square-based pyramidal diamond indenter with included face angles of 136° was penetrated into the material with a force of 20 N. The Vickers hardness number (HV) is related to the applied force and surface area of the resultant permanent impression according to the following equation:

$$ HV = 1.8544 \times \frac{F}{d^2} \quad (5) $$

where F is the force applied and d is the mean diagonal diameter of the impression.

## Scanning Electron Microscopy and Energy Dispersive X-ray Spectroscopy

Microscopic worn surface analysis, generally performed with SEM, furnishes a lot of information about the surface topography, damage, and deterioration level of matrix and filler-matrix adhesion, dominating wear mechanisms, etc. An SEM (Zeiss, EVO-MA 10) was used to study the worn surfaces of the composites. The samples were gold-coated to make them conducting using a Cressington Sputter Coater 108.

## Wettability Analysis

The wettability of matrix and its composites differ based on the type and amount of filler and also filler-matrix bonding. Generally, this property is evaluated by water drop (the sessile drop method), and reduction in the contact angle is an indication of a higher ceramic content and higher filler-matrix adhesion. The contact angle on the surface of selected composites (V, V₂N, and F₂N) was measured using a KRUSS GmbH DSA
<table><caption>**TABLE 3** | Composition and coding of composites.</caption><thead><tr><th colspan="6">(A) Micro-Series with untreated Titania microparticles</th></tr><tr><th>Micro-Series</th><th>TiO<sub>2</sub> (wt.%) Untreated</th><th>0</th><th>2</th><th>4</th><th>6</th><th>8</th></tr><tr><th>Codes</th><th></th><th>V<sub>0</sub></th><th>V<sub>2M</sub></th><th>V<sub>4M</sub></th><th>V<sub>6M</sub></th><th>V<sub>8M</sub></th></tr></thead><tbody><tr><th scope="row">Theoretical density (g/cc)</th><td></td><td>0.930</td><td>0.985</td><td>1.0416</td><td>1.09</td><td>1.153</td></tr><tr><th scope="row">Practical density (g/cc)</th><td></td><td>0.930</td><td>0.952</td><td>0.967</td><td>0.972</td><td>0.996</td></tr><tr><th scope="row">Void content (%)</th><td></td><td>0</td><td>3.3</td><td>7</td><td>10</td><td>13.61</td></tr><tr><th scope="row">Vickers Hardness (MPa)</th><td></td><td>6.621</td><td>6.968</td><td>7.05</td><td>7.15</td><td>7.26</td></tr></tbody></table><table><thead><tr><th colspan="6">(B) Nano-Series with untreated Titania nanoparticles</th></tr><tr><th>Nano-Series</th><th>TiO<sub>2</sub> (wt.%) Untreated</th><th>0</th><th>1</th><th>2</th><th>3</th><th>4</th></tr><tr><th>Codes</th><th></th><th>V<sub>0</sub></th><th>V<sub>1N</sub></th><th>V<sub>2N</sub></th><th>V<sub>3N</sub></th><th>V<sub>4N</sub></th></tr></thead><tbody><tr><th scope="row">Theoretical density (g/cc)</th><td></td><td>0.930</td><td>0.965</td><td>0.98</td><td>1.01</td><td>1.04</td></tr><tr><th scope="row">Practical density (g/cc)</th><td></td><td>0.930</td><td>0.963</td><td>0.967</td><td>0.971</td><td>0.975</td></tr><tr><th scope="row">Void content (%)</th><td></td><td>0</td><td>0.2</td><td>1.32</td><td>3.8</td><td>6.25</td></tr><tr><th scope="row">Vickers Hardness (MPa)</th><td></td><td>6.621</td><td>7.56</td><td>8.14</td><td>9.42</td><td>9.876</td></tr></tbody></table><table><thead><tr><th colspan="6">(C) Nano-Series with treated Titania nanoparticles</th></tr><tr><th>Nano-Series with treated particles</th><th>TiO<sub>2</sub> (wt.%) Treated</th><th>0</th><th>1</th><th>2</th><th>3</th><th>4</th></tr><tr><th>Codes</th><th></th><th>V<sub>0</sub></th><th>F<sub>1N</sub></th><th>F<sub>2N</sub></th><th>F<sub>3N</sub></th><th>F<sub>4N</sub></th></tr></thead><tbody><tr><th scope="row">Theoretical density (g/cc)</th><td></td><td>0.93</td><td>0.963</td><td>0.967</td><td>0.971</td><td>0.975</td></tr><tr><th scope="row">Practical density (g/cc)</th><td></td><td>0.93</td><td>0.964</td><td>0.969</td><td>0.974</td><td>0.979</td></tr><tr><th scope="row">Void content (%)</th><td></td><td>0</td><td>0</td><td>0</td><td>0</td><td>0</td></tr><tr><th scope="row">Vickers Hardness (MPa)</th><td></td><td>6.621</td><td>7.85</td><td>8.21</td><td>9.45</td><td>9.88</td></tr></tbody></table>

25S. The surfaces were cleaned by sonication in acetone prior to the experiment. A 2-µl drop of deionized water was placed on the surface. The contact angles at different positions were averaged out.

## Roughness Measurement

The surface roughness of composites is a tool with which to quantify the texture/topography of the worn surfaces and is correlated with the extent of wear-damage. The surface roughness of the wear track was measured using a Micro-Xam 100 stylus-type profiler manufactured by KLA-Tencor. The roughness data are the average of three points across various areas on the wear track.

## Tribological Characterization

Tribological studies were performed to evaluate the performance of the developed composites in various wear modes, viz., adhesive (sliding against a smooth counterface), abrasive (sliding against a

rougher and harder counterface), and erosive (impingement of harder abrasive particles at different angles) wear modes.

### Adhesive wear

The adhesive wear performance of the composites was evaluated on a UMT-3MT Tribometer supplied by CETR/Bruker, United States. During initial studies on a pin (polymeric) on metallic disk configuration, hardly any loss in weight of the pin was observed, even after running the test for 4 h, because of the exceptionally high wear resistance of UHMWPE. Hence, a ball (AISI 52100 steel – $R_a$ – ≈50 nm) on disk (polymer) configuration was selected to increase the pressure and also the dissipativity of frictional heat, since point contact has these advantages. A 10-mm ball rotated on a polymeric disk in a wear track, as shown in **Figure 3A**. The tests were conducted under 200 N (P 73 MPa), keeping other parameters such as v (0.15 m/s), rpm (100), sliding distance (4220 m), and duration (4 h) constant.
<figure><img src="image_284.png" /></figure>

<figure><img src="image_285.png" /></figure>

<figure><img src="image_286.png" /><figcaption>FIGURE 3 | Configurations in (A) adhesive wear, (B) abrasive wear, and (C) erosive wear.</figcaption></figure>

The specific wear of the disk ($K_0$) (m³/Nm) was calculated as follows.

$$K_0 = \frac{\Delta W}{\rho \times L \times D} (m^3/Nm) \quad (6)$$

where $\Delta W$ – weight loss in kg; $\rho$ – density in kg/cm³, L – load in N; D – sliding distance completed in 4 h in meters.

The maximum Hertzian Pressure ($P$) was calculated as

$$P = \frac{F}{\pi a^2} \quad (7)$$

where $F$ is the applied load (N) and $a$ is the radius of the contact area (mm), calculated as

$$a = \left( \frac{3Fr}{4E^*} \right)^{1/3} \quad (8)$$

where $r$ is the radius of the steel ball and $E^*$ is the calculated modulus at contact.

### Abrasive wear

Abrasive wear studies were conducted in the UMT-3MT Tribometer supplied by CETR/Brucker, United States in a pin on disk (POD) configuration (Figure 3B). Silicon carbide (SiC) abrasive paper of 1000 grade was used to achieve a uniform contact, while for testing purposes, 120 grade (118 µm grit size) was used for a single pass rotary motion. A 10-mm × 10-mm × 4-mm polymer pin was abraded at a constant speed of 0.025 m/s. The sliding distance of 1 m was achieved by changing four papers and was kept constant. In each experiment, all parameters except load were kept constant. Load was selected as a variable parameter and was varied between

25, 50, 75, and 100 N. After completion of the experiment, the pin was cleaned ultrasonically in acetone, dried, and weighed. The specific wear rate was calculated from the recorded weight loss using Eq. 6.

### Erosive wear

The erosive wear tests on the UHMWPE composites were carried out on a TR-471-400 air jet erosion tester (Ducom Instruments Ltd., Bangalore). Ultra-high molecular weight polyethylene,

<figure><img src="image_287.png" /><figcaption>FIGURE 4 | FTIR analysis of siloxane-treated particles for varying treatment durations.</figcaption></figure>
being a ductile material, shows high erosion at lower angles (Hutchings, 1992), and hence 15° and 30° were the two angles chosen for the studies. Alumina particles 60 µm in size were selected as the erodent, and the discharge rate was kept at 4 g/min and the velocity at 58 m/s using a 3-mm diameter Tungsten Carbide nozzle (Figure 3C). The tests were performed at 25°C, and the total mass of erodent was 240 g. Amount of wear was quantified by erosion (E), which was calculated by Eq. 9 (Chahar and Pun, 2018).

$$E = \frac{\text{Weight loss}}{\text{Mass of Erodent}} \text{ g/g} \quad (9)$$

# RESULTS AND DISCUSSION

## Characterization Studies

### FTIR Studies

The FTIR spectra for TiO₂ particles (treated and untreated) are shown in Figure 4.

The broad band observed in the region 600–400 cm⁻¹ corresponded to the -Ti-O-Ti- stretching vibration peak (Ukaji et al., 2007; Mallakpour and Nikkhoo, 2013). In-plane O-H bending vibration corresponds to peaks at 1399 and 1630 cm⁻¹ on the surface of TiO₂. The band at 1088 cm⁻¹ attributed to the bending vibration of Ti-OH is shifted to a lower wavenumber (at 993 cm⁻¹) due to the creation of Ti-O-Si in the spectrum of APTES-TiO₂. The formation of the Ti-O-Si group was confirmed from the FTIR spectrum. Such linkage enables better interaction with the polymeric matrix as compared to the non-functionalized particles. It is expected that the enhanced filler-matrix adhesion would lead to less debonding of particles when stressed during sliding and hence less wear.

## Density

The densities of the composites are given in Table 3. The density of the composite increased with an increase in TiO₂% since it is heavier (4.23 g/cc) than UHMWPE (0.93 g/cc). When comparing the densities of V₂M (0.952 g/cc), V₂N (0.967 g/cc), and F₂N (0.969 g/cc), which contained 2 wt.% Titania particles, it was observed that density was highest where functionalized NPs were used and lowest where untreated MPs were used.

## Vickers Hardness

It can be observed from Table 3 that:

* TiO₂, being a harder filler, increased the micro-hardness of the softer polymeric matrix.

* The hardness of composites increased significantly with increasing content of TiO₂ for MCs and NCs with functionalized particles (F-NCs), indicating good homogeneity of the filler dispersed in the composites. Improvement was highest for F-NCs, ≈60%.

* If the hardnesses of MCs, NCs, and F-NCs with the same amount of particles is compared, they showed the trend V₂M (6.967) < V₂N (8.14) < F₂N (8.20).

# Void Contents

The void contents of the composites showed the following trends.

### F-NCs

* Showed no void contents, indicating ideal composites.

* With increasing contents of MPs and NPs, void contents increased, though not proportionally.

* Void content was highest for MCs followed by NCs, and it was lowest for F-NCs, as follows: V₂M (3.3%) > V₂N (1.32%) > F₂N (0%)

# Wettability Analysis of the Surfaces of Composites

Titania particles, being ceramic, have higher surface energy (72 dynes/cm) than the polymers, UHMWPE (30 dynes/cm) in this case (Travis Curtis, 2013; www.tstar.com, 2020). Hence, it was expected that the inclusion of NPs of Titania in UHMWPE would increase the surface energy of a composite. Their functionalization should do this to a still greater extent. Hence, surface energy or hydrophilicity was evaluated indirectly by wettability studies using a water drop. The UHMWPE showed an angle of around 91°, while those for V₂N and F₂N were 85° and 82°, respectively, indicating increased hydrophilicity.

## Tribo-Properties of Composites

### Adhesive Wear

Adhesive wear tests were conducted initially for 100 N, which resulted in non-measurable wear. Results on adhesive wear at 200 N (73 MPa) are shown in Figure 5. At 300 N, the composites failed at varying sliding durations, mostly due to melting of polymer.

Figure 5A1 shows the effect of inclusion of MPs, NPs, and F-NPs (all 2%) on adhesive wear of UHMWPE. V₀ (virgin UHMWPE) showed the lowest wear rate (5 × 10⁻¹⁵ m³/Nm). Inclusion of Titania particles (irrespective of size and functionalization) led to an increase in friction (Figure 5A2) and wear rate in spite of an increase in the hardness of all of the composites. This was mainly because of the irregular shape of both types of Titania particles (MPs and NPs). Inclusion of particles of solid lubricants such as PTFE and graphite facilitates beneficial film transfer on the counterface, leading to a reduction in metal asperity-contacts and hence reducing friction and wear. Titania particles do not have a layered lattice structure, and hence, beneficial film formation was not expected. An increase in hardness does not ensure an increase in performance in adhesive wear. If the hard particles are spherical, they may act as load-bearing elements and convert area contact to more of a point contact, thereby reducing the probability of metallic asperities penetrating in the polymer composite. This may lead to performance improvement. In the present case, however, irregularly shaped particles with sharp edges posed the severe problem of third-body abrasion. These particles, especially MPs, were debonded easily during initial sliding because of their weaker bonding with the matrix. Once they appear in the debris, they start abrading both the parent surfaces, and adhesive wear is changed to three-body abrasive wear, showing higher friction and wear. MPs were so dominant in third-body
<figure><img src="image_288.png" /><figcaption>A</figcaption></figure>

<figure><img src="image_289.png" /><figcaption>B</figcaption></figure>

<figure><img src="image_290.png" /></figure>

<figure><img src="image_291.png" /><figcaption>C</figcaption></figure>

<figure><img src="image_292.png" /></figure>

**FIGURE 5** | Specific wear rate ($K_0$) (a$_1$,b$_1$,c$_1$) and $\mu$ (a$_2$,b$_2$,c$_2$) of composite disks against a ball (P – 73 MPa, speed – 0.014 m/s, sliding duration – 4 h). (A) Comparison of MC, NC, and F-NC (all with 2% particles) with UHMWPE; (B) Influence of increasing amount of NPs, and (C) benefits of functionalization of NPs.
<figure><img src="image_293.png" /><figcaption>**FIGURE 6** | Failed ball (A) against disks of a microcomposite (B,C) at 300 N and (D) disk of UHMWPE at 300 N.</figcaption></figure>

<figure><img src="image_294.png" /><figcaption>**FIGURE 7** | Worn surfaces of composites disks (A) $V_{2M}$, (B) $V_{2N}$, and (C) $F_{2N}$ tested at 200 N, 0.15 m/s.</figcaption></figure>

<figure><img src="image_295.png" /><figcaption>**FIGURE 8** | Effect of contents of Titania particles on the abrasive wear of UHMWPE as a function of load for (A) microcomposites (MCs), (B) nanocomposites (NCs), and (C) functionalized nanocomposites (F-NCs).</figcaption></figure>

abrasion that during interaction with the steel ball surface, μ was quite high, which led to an excessive increase in frictional torque and hence heat. The polymer disk, having low thermal conductivity (TC), could not dissipate this heat so efficiently, leading to the removal of softened matrix and debonded MPs, easily accounting for the very high wear of this MC (2700% higher than UHMWPE).

In the case of NPs and F-NPs, wear rates increased by 141 and 44% for $V_{2N}$ and $F_{2N}$, respectively. The influence of NPs in composites on performance depends on their type (abrasive or solid lubricant), amount, size, shape, etc. Harder particles, if spherical or cylindrical, can act as ball/roller bearings, reducing

friction and wear in adhesive wear mode, although not much is reported on the influence of the shape of NPs on the adhesive wear of composites (Jia et al., 2006; Shi et al., 2007). In the current case, harder particles were of irregular shape and hence deteriorated the tribo-performance of UHMWPE. For NPs, the digging out process was not so dominant since they had more interaction with the matrix because of their nano size. Hence, debris contained a smaller number of NPs, and third-body abrasion was not so prominent and detrimental. In the case of functionalized NPs, the process of digging out of particles was still more reduced because of stronger filler-matrix adhesion, and the extent of deterioration to the performance also decreased.
<table><caption>TABLE 4 | Average roughness of worn composite disks.</caption><thead><tr><th>Composites</th><th>R<sub>a</sub>(μm)</th><th>Composites</th><th>R<sub>a</sub>(μm)</th><th>Composites</th><th>R<sub>a</sub>(μm)</th></tr></thead><tbody><tr><td>V</td><td>1.45</td><td>1.45</td><td>1.45</td><td>1.45</td><td>1.45</td></tr><tr><td>V<sub>1M</sub></td><td>1.75</td><td>V<sub>1N</sub></td><td>1.54</td><td>F<sub>1N</sub></td><td>1.53</td></tr><tr><td>V<sub>2M</sub></td><td>1.74</td><td>V<sub>2N</sub></td><td>1.76</td><td>F<sub>2N</sub></td><td>1.45</td></tr><tr><td>V<sub>3M</sub></td><td>1.89</td><td>V<sub>3N</sub></td><td>1.77</td><td>F<sub>3N</sub></td><td>1.57</td></tr><tr><td>V<sub>4M</sub></td><td>2.2</td><td>V<sub>4N</sub></td><td>1.84</td><td>F<sub>4N</sub></td><td>1.79</td></tr></tbody></table>

As seen in **Figure 5B1**, an effect whereby an increasing amount of NPs leads to increased friction and wear can be seen, although it is not strictly linear. **Figures 5C1,5C2** show how treatment of particles improved the performance of NCs.

## Failed Surfaces and SEM Studies

Table 4 shows the average roughness ($R_a$) of the wear tracks of composites worn at 200 N. With increasing contents of particles, disk surface roughness increased in almost all composites, which is indirect evidence of third-body abrasion of the disk because of more particles being released in the debris. Interestingly, if $R_a$ values are compared row-wise, they were highest for MPs followed by NPs and lowest for F-NPs, in general, indicating that the severity of third-body abrasion was at a maximum in the case of MPs and at a minimum for F-NPs.

**Figures 6A,D** show photographs of failed ball and disk surfaces of MC when slid under 300 N. The failure was characterized by excessive wear of the disk and adherence of the softened wear debris to the ball surfaces. It also shows localized melting of the wear track at a particular location (**Figure 6C**). An UHMWPE disk (**Figure 6D**) with excessive wear can also be seen.

## SEM of Worn Surfaces of Disks

Although the experiments could run for 4 h, the surfaces show some marks of melting (**Figure 7A**) in the case of MC since frictional heat was very high. The shearing forces on the high-temperature surfaces elongated/fibrillated the UHMWPE. In the case of NC (**Figure 7B**), such marks are negligible (marked), and surface is very rough. In the case of F-NC (**Figure 7C**), fibrillation marks are not evident since the melting phenomenon was milder due to lower frictional heat. However, the entire surface of composite disc shows evidence of softening and then solidifying after the test.

## Abrasive Wear

Incorporation of particles proved beneficial for UHMWPE in the case of abrasion resistance, as seen in **Figure 8** for MCs, NCs, and F-NCs ($K_0$ as a function of contents of Titania particles) and **Figure 9** in comparisons of MCs vs. NCs and NCs vs. F-NCs ($K_0$ as a function of load). It was observed that:

* $K_0$ was in the range of 3–7 × 10⁻¹⁰ m³/Nm for the composites.

* With increase in load, $K_0$ decreased, which is in line with Lhymn's equation (Lhymn, 1987),

$$K_0 = K \frac{v_s}{E H \varepsilon \mu} V_c \frac{1}{F_N^\delta} \quad (10)$$

<figure><img src="image_296.png" /></figure>

where $K_0$ is the specific wear rate, $v_s$ is the sliding speed, $E$ is the elastic modulus, $H$ is the hardness, $\varepsilon$ is the critical wear strain, $\mu$ is the friction coefficient, $V_c$ is the crack growth velocity, and $F_N$ is the normal load. At lower loads, variation was higher. It also shows that wear is inversely proportional to hardness.

* With increase in the amount of MCs (0–8 wt.%), $K_0$ was lowest for 6% followed by a slight rise for 8% (**Figure 8A**). The highest reduction (29%) in $K_0$ was for $V_{6M}$ at 25 N.

* For NCs, the optimum contents for the lowest $K_0$ were in the range of 3%, barring 25 N, where it was 2% (**Figure 8B**), and the highest reduction in $K_0$ was for 2% (22%) at 25N.

* For F-NCs, the optimum contents for the lowest $K_0$ were in the range of 3%, barring 25 N, where it was 2% (**Figure 8C**), and the highest reduction in $K_0$ was for 2% (33%) at 25N.
<figure><img src="image_297.png" /><figcaption>**FIGURE 10** | Worn surfaces of composite pins abraded by SiC paper at 75 N, 25 mm/s for (A) V₀, (B) V₂M, (C) V₂N, and (D) F₂N (secondary micro-cutting indicated by yellow ellipses).</figcaption></figure>

* In general, NCs exhibited superior performance to MCs (Figure 9A).

The reasons for wear resistance enhancement are explained in the section below on worn surface analysis.

## Worn Surface Analysis by SEM

SEM micrographs of composite pins abraded by SiC paper at 75 N are shown in **Figures 10A–D**. The hardness of composites was in the order UHMWPE (6.621 MPa) < V₂M (6.968 MPa) < V₂N (8.14 MPa) < F₂N (8.21 MPa), and wear rate also followed the same order. In abrasive wear, micro-cutting and microplowing are the two major wear mechanisms. Micro-cutting, which is indicated by ribbon-shaped elongated debris and further fibrillation (secondary micro-cutting) due to abrasion by successive grits, and microplowing are the most prominent mechanisms in abrasion for ductile polymers like UHMWPE and PTFE (Hutchings, 1992). Indeed, microplowing occurs in both ductile and brittle polymers. It has a direct correlation with hardness: harder polymers are more resistant to microplowing. In the case of composites with hard fillers like Titania, an additional mechanism of digging out of particles during microplowing also plays an important role. Ease of digging out depends on the size and shape of particles and their bonding with the polymer. Rounded particles can be

easily dug out while being abraded by the SiC grits, while irregularly shaped particles resist to a greater extent. After being dug out, the particles appear in wear debris. These then start damaging the SiC grits and their abrasivity. The higher the hardness, the higher is the damaging power, and the higher is the protection to the polymer composite from the grits during further abrasion.

As seen in Micrograph 10a, excessive micro-cutting and secondary micro-cutting were responsible for the highest wear of UHMWPE. The depth of grooves due to microplowing is also highest, since UHMWPE has the lowest hardness. Moving from UHMWPE to V₂M, hardness increased, and both microplowing and micro-cutting were abruptly reduced. Drastically decreased ductility of the composite can be seen from the edges of ribbons (marked). The depth of grooves is also reduced, indicating reduced extents of both phenomena, supporting the lower wear of this MC.

For NC (Figure 10C) and then F-NC (Figure 10D), increased ductility (due to their being NCs) can be observed along with secondary micro-cutting. The microplowing is drastically reduced, especially for F-NC, where it is reduced to the extent that the surface topography appears to be quite smooth. This supports the lowest wear occurring in this composite.

**Supplementary Appendix A** elaborates on the difference in adhesive and abrasive wear and the reasons behind it.
<figure><img src="image_298.png" /><figcaption>**FIGURE 11** | Erosion as a function of amount of particles for two angles; (A) microcomposites (MCs), (B) nanocomposites (NCs), and (C) functionalized nanocomposites (F-NCs); comparative erosion performance for (D) MCs and NCs, and (E) NCs and F-NCs.</figcaption></figure>

## Erosive Wear

Erosive wear is a complex wear mode, and in case of particulate composites, wear resistance is always deteriorated, mainly due to the ease with which particles are detached and thrown off. The literature reports that in spite of increasing hardness, wear performance always decreases (Suh, 1986). An enhancement of performance has only been reported for Aramid and Zylon fabric reinforced composites (Sharma et al., 2019). Figure 11 shows the erosion performance of composites eroded at two angles. In all cases, erosion of composites was higher at 15°, indicating super-ductile failure (Hutchings, 1992).

It was observed that the virgin UHMWPE was the best performer. Inclusion of Titania particles deteriorated the performance significantly in spite of an increase in hardness. The extent of deterioration increased with increase in Titania particle content, being at a maximum for MPs and a minimum for F-NPs. Erosion definitely depends on the hardness of the analyte material. In the case of composite, which is a heterogeneous material, however, hardness does not help. It depends on the ease with which particles are dislodged from the matrix during bombardment. NPs definitely lead to more homogeneity of

the composite and have many-times superior particle-matrix adhesion because of the increased surface activity of NPs. It is not so easy to dislodge NPs, and hence wear of NCs was significantly less severe than that of MCs. Interestingly NC with 1% Titania showed quite similar behavior to UHMWPE. For functionalized NCs, adhesion was still stronger, and hence erosion performance was the best among all of the particulate composites. Figures 11D,E demonstrates this comparative behavior more clearly.

## CONCLUSION

In this study, composites of UHMWPE and Titania were developed with three objectives, viz, to investigate the size effect (nano vs. micro), the amount effect, and the particle-functionalization effect.

It was worth pointing out that the testing conditions in this work were not in tune with testing for biomedical applications, although UHMWPE composites are being researched for the same. There is scope to extend this work in that direction. This
work is instead an indication of how abrasive wear resistance of UHMWPE can be increased by incorporating functionalized Titania NPs of irregular shape. The same theme may possibly work to enhance adhesive wear resistance if spherical NPs with functionalization could be used.

Based on the investigations on the three series of composites, the following conclusions are drawn.

* Functionalization of NPs led to increased adhesion with the matrix, as evident from the increased density (compactness) and lowering of contact angle (increased hydrophilicity).

* Inclusion of particles increased the hardness of composites. The higher the amount, the higher was the hardness. For composites with the same amount of particles (2%), functionalized nanocomposite (F-NC) showed the highest hardness, followed by NC, and then MC.

* Based on the abrasive wear studies, the optimum% of MPs for the highest improvement (20%) in wear resistance was 6%, while for nano-series, it was 3% for attaining a 28% improvement. Functionalized NPs proved most effective, and a 35% improvement in wear resistance was observed for 3 wt% of F-NPs. Increase in hardness was found to be the major reason for this, as it was responsible for reducing the severity of dominating wear mechanisms viz micro-cutting and microploving significantly, as confirmed from the SEM studies.

* In the case of adhesive and erosive wear modes, however, the inclusion of particles led to a deterioration in performance, which depended on the size, amount, and functionalization of NPs. It was concluded that the increase in hardness did not help to increase the wear performance in these two wear modes. The irregular shape of MPs and NPs was the main controlling parameter. Including 2% MPs deteriorated the performance of UHMWPE excessively (2400%) in adhesive wear mode, while NPs reduced it by 140%. F-NCs were least detrimental (40% deterioration). Third-body abrasion appeared to be the dominating wear mechanism due to released particles from the composites accumulating as wear debris in adhesive wear mode. An increase in the roughness of worn surfaces supported this phenomenon.

* In the case of erosive wear, UHMWPE again proved best, followed by F-NCs and then NCs. MCs proved worst, and the higher the amount of particles, the more severe was the deterioration. It was concluded that the ease with which particles are dislodged from the composites during bombardment was the key factor controlling the erosion

of composites. F-NCs proved best in this respect since particle-matrix bonding was the best. The composites showed failure in super-ductile mode by exhibiting the highest erosion at 15°.

The most important conclusions regarding the selection of particles to improve the hardness and wear resistance of UHMWPE are as follows.

The type, size, amount, shape, and filler-matrix adhesion are the most important aspects. Nano-sized particles worked best, and a content of around 3% proved optimum. Siloxane treatment proved very effective in improving all performance properties. However, the shape of particles controlled adhesive wear performance in an adverse way. For abrasive wear mode, it was appropriate.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

## AUTHOR CONTRIBUTIONS

MP conducted and repeated many experiments and performed the manuscript writing and data analysis. GP developed composites, conducted several experiments, and helped in data compilation along with figures and tables. JB performed to identifying and conceptualizing of the problem and supervision during implementation, analysis of the data, helped in writing the draft of the manuscript, and contributed extensively for the depiction of the ideas for the present. All authors contributed to the article and approved the submitted version.

## ACKNOWLEDGMENTS

The authors are thankful to Mr. Bhaskaranand Bhatt (Research Scholar, Indian Institute of Technology Delhi) for providing help in SEM work.

## SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats.2020.00205/full#supplementary-material

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