The effect of resin properties on toughness translation in interleaved composites

Abstract

Interleaving is a well-recognized method for enhancing the interlaminar fracture toughness of composite materials. However, theoretical toughness is often not realized since cracks tend to propagate through regions of lower toughness. While most studies have focused on the effect of interleaf properties, e.g., thickness and interleaf resin properties, on composite interlaminar fracture toughness, the effect of matrix resin properties on interleaved composite toughness has been overlooked. Recently, we hypothesized that there is a relationship between toughness translation and the ratio of matrix to interleaf resin toughness. In this work, we use additive manufacturing to test this hypothesis with a range of resins in interleaved composite. Toughness is quantified via mode I delamination resistance. Our results confirm our hypothesis that the ratio of matrix to interleaf resin fracture toughness, i.e., the ratio of $K_{I C}$ , is directly correlated to the degree of toughness translation. More specifically, a ratio of approximately 0.5 is required for effective interleaving; below 0.5, the toughness translation is significantly below the theoretical value. This approach has important implications for the design and fabrication of tough composites via choice of resins.

Keywords

Interleaved composite toughened composites multi-resin composites interleaving additive manufacturing resin rich layers

Introduction

Laminated fiber reinforced composites are multilayered structures with exceptional strength and stiffness. However, their modest interlaminar toughness and susceptibility to delamination failure present significant challenges.^1–3 Interlaminar toughening offers a promising solution to these limitations, enabling the broader utilization of composites in high-end applications.^4–6 Currently, three major approaches to interlaminar toughening exist: through-thickness reinforcement (z-pinning), matrix toughening, and interlayer toughening (interleaving).¹ Each technique provides a distinct level of toughness improvement, but also comes with specific drawbacks that require careful consideration, namely, possible disturbance of fiber architecture, reduction of fiber volume fraction, and loss of mechanical properties.^7,8

Matrix toughening is one of the first methods employed to reduce delamination failure. This method has the least effect on fiber volume fraction for matrices with reasonable viscosity. Previously, it was thought that incorporating a tough matrix would directly enhance interlaminar toughness. However, subsequent studies revealed that highly toughened matrices are not as effective in toughening the composite due to the restriction of the plastic deformation zone.^9–11 In other words, the interlaminar region size significantly affects resistance to crack growth. The region must be comparable to the plastic deformation zone to achieve significant toughening of the composite. The process of introducing a thick resin rich layer (RRL) into the laminate to allow for plastic deformation and suppress crack growth is known as interleaving and is effective in improving both mode I and mode II interlaminar toughness.

The first-generation interleaf materials consisted of discrete thermoplastic or thermoset films placed between laminate layers.¹² However, incorporating RRL/interleafs in composite materials is not trivial using standard manufacturing techniques. One solution was to use fillers, such as fiber veils, particles, short fibers, and nano materials like graphene and carbon nanotubes, as spacers to control the interlayer thickness and expand the range of interleaf materials.^13–16 Although these fillers offer processing advantages, they do not necessarily participate in the toughening phenomena.^17,18 Therefore, it is not clear wether fillers are need to be used in interleaving methods to achieve toughened composite materials. Furthermore, the use of fillers still limits the composite to a single resin.

Additive manufacturing offers a strategic platform for directly incorporating RRLs into composite parts and using multiple resins for laminate and interleaf.¹⁹ In this technique, no fillers are required and the RRL thickness and chemistry can be prescribed accurately with a z-axis motor and different resin vats.¹⁹

Under certain circumstances, the theoretical performance of interleaved composites can be predicted using Linear Elastic Fracture Mechanics (LEFM). For instance, when the interleaf is thick enough, i.e. on the length scale of the plastic zone size, and there is no other phenomenon such as fiber bridging, fracture of the composite satisfies plane strain conditions such that a cohesive failure is expected, and a composite interlaminar fracture toughness ( $G_{I C}$ ) equivalent to the resin-rich layer (RRL) toughness ( $G_{I C}$ ) is observed.^20,21 This is shown in previous studies on self-interleaved composites and in the performance of recently reported hybrid toughened and veil toughened composites.^15,20,22 However, these results are all restricted to single resin properties across the entire composite part.

In heterogeneous systems where the interleaf material differs from the matrix, adhesive failure is commonly observed where cracks tend to propagate in the brittle matrix rather than cohesively through the tough interleaf.^7,13,17,23 Despite reports of toughness improvement, such failure excludes the interleaf from the fracture process at an early stage and diminishes the improvement window. Intuitively, one can assume this occurs when the interleaf is stronger than the matrix, as the crack would follow the path of least resistance, i.e., the brittle matrix. We recently hypothesized that the ratio between the resin toughness in the lamina and the RRL is a key factor in achieving cohesive failure.¹⁹ This argument is supported by Ni et al., who demonstrated through finite element simulation that if the interlayer bulk toughness ( $G_{I C} a n d G_{I I C}$ ) are only 10% higher than that of the base lamina, failure would be adhesive, i.e. the crack propogates along the fiber matrix interface.²⁴ This hypothesis, if validated, has important implications on the design of toughened interleaved composites.

In this work, we used a range of resins with incremental differences in fracture toughness to verify that there exists an ideal ratio of resin toughness for realizing cohesive failure and maximizing interlaminar toughness translation in interleaved composites. We utilized vat polymerization additive manufacturing, a previously validated method, to produce laminated composites with controlled RRL thickness and tested mode I interlaminar fracture toughness of control and interleaved composites with different combinations of resin. This technique is discussed in detail and compared with vacuum-assisted resin transfer molding in Idrees et al.²⁵ Note that controlled interleaved specimens could not be made without the use of additive manufacturing.²⁵ The results obtained are combined with findings from a previous study to establish a correlation between resins’ fracture toughness ratio, composite mode I fracture mechanisms, and overall toughness translation.¹⁹

Materials and methods

Materials

DA2 is a stiff resin consisting of bisphenol A glycerolate dimethacrylate (Bis-GMA), ethoxylated bisphenol A dimethacrylate (Bis-EMA), and 1,6-hexanediol dimethacrylate (HDDMA) mixed at a 3:3:2 weight ratio, along with 0.7% wt of phenylbis (2,4,6 trimethylbenzoyl) phosphine oxide photoinitiator (PPO). DA2 mechanical and thermal properties were published by Tu et al.²⁶ DA5 is an in-house developed vinyl ester resin synthesized by reacting two types of epoxy resin, namely, Epon 828 and Epon 1001F, with methacrylic acid to produce methacrylated epoxy resin. Para-methyl styrene is added as a reactive diluent. The resin-diluent ratio is 65:35, respectively, and 0.7 wt% PPO is used as a photoinitiator. The formulation and properties of DA5 were published by McLaughlin and Palmese.²⁷ Additionally, a urethane acrylate commercial resin with the brand name “Tenacious” was purchased from Siraya Tech. The mechanical properties of Tenacious are published in reference.^19,28 The DA2 and Tenacious were mixed at different weight ratios to formulate resins with varying fracture toughness, referred to here as DT-XX, where XX is a number representing the weight fraction of Tenacious in the mixture (∼40%–80%). DT-XX resins are formulated by mixing the two resins for one hour. Details about their formulation and mechanical and fracture properties were published in reference²⁸ A fabric, namely, plain weave E-glass with an areal density of 0.025 g/cm², thickness of 0.025 cm and a single fiber diameter of ∼12.5 μm, was purchased from FiberGlast Developments Corp. The fabric contains general-purpose sizing.

Composite manufacturing

Figure 1 shows composite processing via vat polymerization additive manufacturing, using a process similar to what was previously described in reference.¹⁹ We utilized an ELEGOO MARS LCD (Liquid Crystal Display) printer equipped with a 405 nm LED to fabricate laminated composites using a modified technique that involved applying pressure on fiber mat during the printing process. This method allowed us to interchange the resin vat to incorporate different matrices, enabling the creation of resin-rich layers (RRLs). To cure the composites, we consolidated 4 layers of fiber mats at a time using the building platform and applied a 150 N of force. Subsequently, the cured fiber mats underwent exposure to a defined pattern of 405 nm light for a specified time. Our composite parts were made of 16 fiber mat layers, so the part was made in four steps; during printing, three rectangular specimens were produced simultaneously, with separate boundaries for easy handling. We swapped the resin vat to accommodate different resin types, specifically for the RRL printing. For RRLs printing, we employed a slicing thickness of 50 µm and an exposure time of 100 seconds, resulting in a ∼200 µm RRL. After printing, excess resin was wiped with a kimwipe saturated with isopropanol. The specimen was then dried and photo post-cured at 75°C for two hours. Finally, specimens were polished using a polishing wheel (Spectrum System 1000 polisher) and a series of coarse to fine grit sandpaper. The RRL thickness was measured using an optical microscope. Pre-cracked specimens were checked with a microscope to ensure that the crack tip was near the midplane of the RRL. Specimens with crack tips significantly off centered were observed but not tested. Table 1 provides additional details regarding the fabricated composites.

Figure 1.

(a) Schematics of the AM process used for fabrication of laminated composites with RRL, and (b) schematics of a DCB test specimen and our research question on how RRL and matrix fracture toughness translate to the final composite fracture toughness.

Table 1.

Summary of printed composite specimens with their reference names, materials, and resins fracture toughness.

	Specimen code	Fiber resin (matrix)	RRL resin (interleaf)	Average RRL thickness (µm)	Total specimen thickness (mm)	Specimen width (mm)	Matrix $K_{I C}$ (MPa.m^0.5)	RRL $K_{I C}$ (MPa.m^0.5)	RRL-matrix $K_{I C}$ ratio ( $Φ_{k})$	Ref
1	DA2-GF	DA2	NA	NA	-	-	0.45 ± 0.02 [26]	-	-	²⁵
2	DT40-GF	DT40	NA	NA	3.35 ± 0.05	19.86 ± 0.53	0.55 ± 0.08 [28]	-	-	This work
3	DT60-GF	DT60	NA	NA	3.36 ± 0.03	19.83 ± 0.35	0.75 ± 0.09 [28]	-	-	This work
4	DT80-GF	DT80	NA	NA	3.37 ± 0.03	19.80 ± 0.23	1.22 ± 0.12[28]	-	-	This work
5	Tenacious-GF	Tenacious	NA	NA	3.45 ± 0.09	20.25 ± 0.23	1.8 ± 0.26[28]	-	-	This work
6	DA2-GF-tenacious-200	DA2	Tenacious	206±6	3.75 ± 0.13	19.82 ± 0.12	0.45 ± 0.02[26]	1.8 ± 0.26[28]	0.25	¹⁹
7	DT40-GF-tenacious-200	DT40	Tenacious	205 ± 15	3.56 ± 0.04	19.26 ± 0.59	0.55 ± 0.08[28]	1.8 ± 0.26[28]	0.30	This work
8	DT60-GF-tenacious-200	D60	Tenacious	227 ± 32	3.49 ± 0.04	19.38 ± 0.4	0.75 ± 0.09[28]	1.8 ± 0.26[28]	0.41	This work
9	DT80-GF-tenacious-200	DT80	Tenacious	209 ± 18	3.57 ± 0.06	19.37 ± 0.46	1.22 ± 0.12[28]	1.8 ± 0.26 [28]	0.67	This work
10	DA2-GF-DA5-200^a	DA2	DA5	201 ± 30	3.57 ± 0.01	19.57 ± 0.21	0.45 ± 0.02[26]	0.6 ± 0.07 [27]	0.75	¹⁹
11	DA5-GF-tenacious-200^a	DA5	^aTenacious¹	219 ± 20	3.72 ± 0.06	19.63 ± 0.23	0.6 ± 0.07[27]	1.37 ± 0.06[19]	0.43	¹⁹
12	DA2-GF-DA2-200	DA2	DA2	199 ± 21	3.52 ± 0.01	18.95 ± 0.23	0.45 ± 0.02[26]	0.45 ± 0.02[26]	1.0	¹⁹

^aThis composite went through a second stage of thermal post cure at 120°C.

Mode I interlaminar fracture toughness, resin toughness, and resin fracture toughness

Mode I interlaminar fracture toughness tests were conducted according to ASTM D5528 standards using an Instron apparatus with a 1.0 kN load cell. Double cantilever beam (DCB) specimens, measuring ∼20 mm in width and 120 mm in length, were printed. An RRL of ∼200 µm thickness was incorporated into the interleaved specimens’ midplane. A precrack of approximately 50 mm in length was created using a 12.7 µm thick polyimide film. Aluminum piano hinges were attached to the specimens with a superglue adhesive, and the sides were marked for tracking crack length over time. Five specimens of each set were tested. The DCB test was conducted at room temperature and crosshead speed of 5.0 mm/min. The crack length was monitored using a video camera. When possible, the interlaminar fracture toughness parameter $G_{I C}$ , also referred to as the strain energy release rate per area, was calculated using the modified beam theory method (MBT)

G_{IC} = \frac{3 P δ}{2 b (a + | Δ |)}

(1)

where

P

is the measured force,

δ

is the measured extension,

b

is the specimen width,

a

is the measured crack length, and

Δ

is the effective delamination extension to correct the DCB arms’ rotation.

G_{I C}^{i n i t}

is the measured

G_{IC}

value at a = 50 mm, while

G_{I C}^{s s}

is the average of all

G_{IC}

value for a≥ 60 mm. Note that, in any instances, a true steady state is not reached. Note that the G_IC of pure resins can be measured using the same Mode I test, however, the resulting G_IC is referred to as simply the resin toughness, since there is no interlaminar regions. While this is related to the plane strain fracture toughness, K_IC, the analysis is different. All K_IC values reported in this work were measured previously using ASTM D 5045 and are reported in Table 1 with their corresponding reference.^19,26–28 Similar to reference,^9,10,29,30 and ASTM D5045, in this manuscript, pure resin G_IC, i.e. without laminates, is referred to as resin toughness and K_IC is used to refer to resin fracture toughness.

Results and discussion

Control composites mode I interlaminar fracture toughness

It is well established that mode I delamination failure occurs at lower energy levels than mode II and is influenced by the resin toughness ( $G_{I C}$ ) and fiber-resin interface strength.^9,31,32 Previous research established a correlation between tensile yield stress and or matrix toughness ( $G_{I C}$ ), and composite interlaminar fracture toughness.^10,21,33,34 In brittle matrices, $G_{I C}$ is typically higher than or equal to the matrix toughness, with the fiber-matrix interface playing a significant role in increasing toughness. However, when the matrix toughness exceeds a certain level, its toughness is limited by the restricted resin deformation zone.^11,33,35

The DCB test was used to investigate mode I interlaminar fracture toughness, and the resulting $G_{I C}$ values were compared to the matrix toughness. Figure 2 shows the test results for mode I interlaminar fracture toughness with no RRL, i.e., the control specimen. Figure 2(a) shows the DCB load displacement curves of the four composites. The Tenacious glass fiber composite showed a peak load of approximately 50 N, followed by a gradual decline due to crack propagation. The stable crack propagation in Tenacious composites was attributed to significant fiber bridging, which is evident from images taken from the DCB test (Figure 3). The only other composite that showed a stable crack propagation was DA2-GF. The other resin composites (DT40-GF, DT60-GF, and DT80-GF) exhibited periodic sharp load drops resembling saw-teeth, which is often called stick-slip behavior.^36–38 The magnitude of these drops increased with the concentration of Tenacious resin in the blend. These large load drops are not typical of mode I failure as predicted by ASTM standards, which anticipate stable growth and small crack length increments.

Figure 2.

DCB test results of control laminates made with resins of varying toughness(a) DCB load displacement curves (b) R-curves, toughness vs. crack length (c) Laminates average initiation and steady state $G_{I C}$ vs. the matrices toughness(d) Top view of different DCB specimen fracture surfaces.

Figure 3.

Snapshot of Tenacious-GF composite DCB test showing fiber bridging.

The load drops and unstable crack growth are reflected in the R-curves (see Figure 2(b)), where toughness is plotted as a function of crack length. Figure 2(b) shows fluctuations in toughness between an upper and lower bound, while DA2 and Tenacious reached a nearly steady state. In the case of DT80, $G_{I C}$ fluctuated between ∼0.9 and 1.5 kJ/m².

In Figure 2(c), we show the effect of matrix toughness on $G_{I C}$ at initiation, $G_{I C}^{i n i t}$ , and steady-state toughness, Except for the case of Tenacious composites, it is evident that both initiation and steady-state toughness increased with increasing matrix toughness and were higher than the matrix toughness. $G_{I C}^{s s}$ was higher than $G_{I C}^{i n i t}$ , which is typically attributed to fiber bridging.¹⁰ On the contrary, the $G_{I C}^{i n i t}$ of Tenacious composites are lower than the matrix $G_{I C}$ . This is common in resins with high toughness, where the interlayer thickness was significantly lower than the matrix plastic zone size.³⁵ Thus, these results agree qualitatively with literature findings that beyond a specific matrix toughness, an increase in the matrix toughness does not translate to the composite. Note that for very tough resins, the composite toughness can be as low as one-third of the resin toughness.^11,35

The periodic interlaminar toughness observed in Figure 2 can be due to a range of factors. One explanation is a change in either the fiber-matrix interface properties or the matrix bulk properties. A study on vinyl ester and glass satin woven fabric showed that stick-slip behavior (unstable crack growth) increased with the concentration of the sizing agent, indicating that the interface was the primary contributor to this phenomenon.³⁹ Another explanation is the matrix’s sensitivity to the loading rate. A study on graphite PEEK (Polyether ether ketone) demonstrated that stick-slip behavior was due to the ductile-to-brittle transition of the resin at different deformation rates. At low rates, stable crack growth was observed, while at high rates, unstable crack growth occurred when the cracks reached the boundaries.³⁶

The fracture surface of the DCB specimens (Figure 2(d)) showed signs of stick-slip behavior in the blended resins. DA2 composite failure exhibited clear adhesive failure, while blended resin composites showed significant resin presence on both fractured surfaces. The white regions in Figure 2(d), particularly for DT60 and DT80, correspond to the stick (crack arrest) region, where an increase in interlaminar toughness is observed, followed by long regions of cohesive failure. The length of unstable crack growth and the process zone increased with increasing resin toughness. This could be due to ductile-to-brittle transitions since the interfacial properties are not expected to differ greatly. Overall, even though Tenacious composites had the highest toughness, blending DA2 and Tenacious improved composite interlaminar toughness over the brittle DA2 resin composite.

Mode I interlaminar fracture toughness of interleaved composites

In this part, we examine the impact of matrix toughness on interleaved composite toughness, specifically focusing on the effect of Tenacious RRL (Resin Rich Layer) with a thickness of approximately 200 microns. The RRL is expected to enhance the interlaminar toughness of the composite, bringing it closer to the bulk toughness of the interleaf material. Theoretically, the interleaved composite toughness should equal the RRL resin toughness (1.5 kJ/m²) when the thickness is greater than the plastic deformation zone.^19,20

Figure 4 shows the DCB test results for Tenacious interleaved composites. Figure 4(a) shows representative load-displacement curves for each specimen. Note that in all cases, the peak load is higher than the non-interleaved composites discussed in Figure 2. However, we still observe stick-slip behavior for certain resins, as evidenced by significant load drops. The improvement in fracture toughness is better seen in Figure 4(b). In all cases, the composite initiation toughness exceeds the steady state toughness, highlighting the role of the RRL in resisting crack initiation. Although the toughness decreases after initiation, the steady state $G_{I C}$ remains significantly higher than those of the control sample in Figure 2.

Figure 4.

Interleaved composites DCB results (a) DCB load-displacement curve, (b) R-curves, toughness as a function of crack length, (c) the effect of matrices toughness on interleaved composites performance, (d) Top view of DCB fracture surfaces of interleaved composites.

Figure 4(c) summarizes the average initiation and steady state toughness of the interleaved composites and compares them to matrix toughness. Unexpectedly, we observe that both initiation and steady state toughness improve linearly with increasing matrix toughness. Recall that we expected the fracture toughness to depend purely on the interleaf resin properties, which are constant for all specimens. This result is followed by unexpected differences in the fracture surface (Figure 4(d)), where the failure is predominantly adhesive for DA2, DT40, and DT60 interleaved composites. In the case of DT80, only a small adhesive zone (process zone) is observed near the insert tip followed by a fully planar cohesive failure, which is qualitiately consistent with literature.^19,24However, unlike the numerical simulation of Ni et al. which shows that a ratio of G_IC (interlayer) to G_IC (control laminate) greater than 10% leads to a transition from cohesive to adhesive failure, the experimental results show that this transition occurs at 180% difference in the G_IC ratio.^11,24,35

These findings clearly demonstrate that the failure and toughness of interleaved composites are strongly dependent on the matrix properties, which is not an expected result. In a previous study, we established that a minimum ratio of interleaf to matrix resin $G_{I C}$ equal to 7.2 was necessary for achieving full toughness translation, i.e., matching or exceeding the toughness of the interleaf resin. Here, we propose using the ratio of plane strain fracture toughness $K_{I C}$ of the resins, which is a stress based measure of resin toughness, instead of the ratio of resin $G_{I C}$ . This choice leads to a linear trend with toughness translation. We assume that crack initiation within the interleaf or the matrix is determined by the relative fracture toughness of both resin materials. Specifically, toughness translation is directly proportional to the matrix $K_{I C}$ and inversely proportional to interleaf $K_{I C}$ . We define toughness translation, $η$ , as the ratio of the interleaved composite initiation toughness $G_{I C}^{i n i t}$ and the $G_{I C}$ of the RRL resin. $η = \frac{G_{I C} (i n t e r l e a v e d c o m p o s i t e)}{G_{I C} (R R L r e s i n)}$ .

Figure 5 relates toughness translation to individual $K_{I C}$ , and the ratio of matrix to interleaf $K_{I C}$ , $ϕ$ , defined as

ϕ = \frac{K_{I C} (M a t r i x r e s i n)}{K_{I C} (R R L R e s i n)} .

Figure 5.

The effect of RRL and matrix properties on toughness translation, (a) the effect of matrix fracture toughness on toughness translation, (b) the effect of RRL resin fracture toughness on toughness translation, (c) The combined effect of both RRL and matrix fracture toughness on toughness translation.

Figure 5(a) illustrates the influence of matrix $K_{I C}$ on η. Interestingly, there is no unique dependence of η on matrix $K_{I C}$ . More specifically, while there is a general increase of $η$ with matrix $K_{I C}$ , we observe a range of $η$ values when the interleaf $K_{I C}$ is varied and matrix resin $K_{I C}$ is constant (indicated by the green arrows). Similarly, Figure 5(b) shows a general decreasing trend of $η$ with increasing interleaf $K_{I C}$ , but $η$ varies strongly with matrix $K_{I C}$ when interleaf $K_{I C}$ is kept constant. Only when $η$ is plotted as a function of $ϕ$ a unique trend is observed. Figure 5(c) shows a monotonic increase in toughness translation with increasing values of $ϕ$ , followed by an asymptotic value. We observe a region of poor translation for $η <$ 1. Although we expected a plateau at $η =$ 1, we observe a plateau at 1.8, which could be attributed to variation in notch sensitivity between SENB (Single Edge Notched Bending) and DCB tests, as well as possible fiber contributions in the initiation. For $ϕ$ > 0.5, we observe excellent toughness translation, above the theoretical limit, while complete toughness translation is observed for $ϕ \geq 0.7$ . More work is needed to understand whether these values of $ϕ$ are restricted to the chemistry measured in this work. Regardless, these results clearly emphasize the significance of the relative properties of the interleaf and matrix in attaining high interlaminar toughness.

Conclusion

In this study, we utilized additive manufacturing to produce heterogeneous interleaved composites that could not be made using traditional manufacturing techniques. The AM method was used to control the RRL thickness and the resin chemistry for both the matrix and RRL. These specimens were used to determine a relationship between resin fracture toughness and composite interlaminar toughness. Firstly, we observed that both interleaved and non-interleaved composites exhibited increased initiation fracture toughness as the matrix toughness was enhanced. This indicates that the matrix resin plays a crucial role in determining the overall toughness of heterogeneous interleaved composites.

The results clearly indicate a correlation between the relative toughness of the matrix and the effectiveness of interleaving. More specifically, there exists a minimum fracture toughness ratio ( $ϕ$ ∼0.5) to observe full theoretical translation, while $ϕ \geq$ 0.7 is necessary to obtain a maximum translation. Overall, these results provide valuable guidelines for the selection of materials in laminates aimed at enhancing delamination resistance. More work is needed to determine if the restrictions on $ϕ$ are the same for other resin systems and different fiber architectures.

Footnotes

Acknowledgments

This research was sponsored by the Army Research Laboratory and was accomplished under Cooperative Agreement Number W911NF-17-2-0227. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of the Army Research Laboratory or the U.S. Government. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes not withstanding any copyright notation herein.

Declaration of conflicting interests

The author(s) declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Funding

The author(s) disclosed receipt of the following financial support for the research, authorship, and/or publication of this article: This work was supported by the Army Research Laboratory (W911NF-17-2-0227).

ORCID iD

Nicolas J. Alvarez

Data Availability Statement

Data is available upon request.*

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