Z2 and Z3 ones, which, in turn, explains the reason why the prepreg thickness at Z2 and Z3 zones is higher than the one in the adjacent ones. At increasing holding force, the interfacial pressure between the FML and dies is higher, which induces a lower prepreg thickness.

<figure><img src="image_460.png" /><figcaption>Fig. 4 Prepreg thickness distribution.</figcaption></figure>

Fig. 5 shows some details of the thickness distribution at various zones when applying the highest blank-holding force. The thickness of the AZ31B sheets (upper and lower layers) shows a uniform distribution while a non-uniform thickness distribution of the prepreg can be observed at different zones. The thickness of the prepreg at Z1, Z4, and Z5 zones is much lower than that at Z2 and Z3 zones.

<figure><img src="image_461.png" /><figcaption>Fig. 5 Detail of the FML thickness distribution when applying the highest blank-holding force.</figcaption></figure>

Forming force. Fig. 6 reports the forming force at varying blank-holding forces. The force gradually increases, till it reaches a stable level. Before # 2, the FML bends at Z5 zone and the punch only touches the Z4 zone: as a result, the higher contacting pressure at Z4 zone squeezes the prepreg to the adjacent areas facing lower normal pressure, therefore leading to the prepreg thinning at Z4 zone compared to the adjacent zones (see the experimental results in Fig. 7 (a) under the lowest blank-holding force at # 1). As the punch moves down, the Z5 zone of the FML touches the lower die at # 2, as shown in Fig. 7 (b). From # 2 on, the forming force increases dramatically to flatten the bent FMLs at Z5 zone; in the meanwhile, a higher normal pressure is induced and the prepreg at this area is squeezed to both the adjacent areas and outside, resulting in a decrease in the thickness at Z5 zone, as shown in Fig. 7 (c) at # 3.
Z2 and Z3 ones, which, in turn, explains the reason why the prepreg thickness at Z2 and Z3 zones is higher than the one in the adjacent ones. At increasing holding force, the interfacial pressure between the FML and dies is higher, which induces a lower prepreg thickness.

<figure><img src="image_471.png" /><figcaption>Fig. 4 Prepreg thickness distribution.</figcaption></figure>

Fig. 5 shows some details of the thickness distribution at various zones when applying the highest blank-holding force. The thickness of the AZ31B sheets (upper and lower layers) shows a uniform distribution while a non-uniform thickness distribution of the prepreg can be observed at different zones. The thickness of the prepreg at Z1, Z4, and Z5 zones is much lower than that at Z2 and Z3 zones.

<figure><img src="image_472.png" /><figcaption>Fig. 5 Detail of the FML thickness distribution when applying the highest blank-holding force.</figcaption></figure>

Forming force. Fig. 6 reports the forming force at varying blank-holding forces. The force gradually increases, till it reaches a stable level. Before # 2, the FML bends at Z5 zone and the punch only touches the Z4 zone: as a result, the higher contacting pressure at Z4 zone squeezes the prepreg to the adjacent areas facing lower normal pressure, therefore leading to the prepreg thinning at Z4 zone compared to the adjacent zones (see the experimental results in Fig. 7 (a) under the lowest blank-holding force at # 1). As the punch moves down, the Z5 zone of the FML touches the lower die at # 2, as shown in Fig. 7 (b). From # 2 on, the forming force increases dramatically to flatten the bent FMLs at Z5 zone; in the meanwhile, a higher normal pressure is induced and the prepreg at this area is squeezed to both the adjacent areas and outside, resulting in a decrease in the thickness at Z5 zone, as shown in Fig. 7 (c) at # 3.
Before # 2, in order to overcome the higher friction force between the FML and dies due to the higher blank-holding force, a higher forming force is required. While the forming forces under Hf₁ and Hf₂ are almost the same after # 2, meaning that there is almost no relative sliding during the flattening process of the Z5 zone.

<figure><img src="image_462.png" /><figcaption>Fig. 6 Forming forces under different holding forces</figcaption></figure>

<figure><img src="image_463.png" /><figcaption>Fig. 7 Thickness distribution at # 1(a) 2 (b) and 3 (c) under the lowest blank-holding force.</figcaption></figure>

**Summary**

In this study, the thermoforming process was conducted on thermoplastic polymer and magnesium alloy-based FMLs to manufacture hat-shaped components. The effect of the blank-holding force on the forming results was analysed by comparing the forming force during the tests and the thickness distribution of each layer of the FML. The key findings can be summarized as follows:

* A non-uniform thickness distribution was found at different areas of the hat-shaped FML parts, and a significant reduction of the prepreg thickness was observed at the bottom radii compared to other areas.

* Moreover, the higher the blank-holding force the higher the forming force and the more significant the prepreg thickness variation.

**Acknowledgments**

This research was supported by Ministry University Research within the PRIN 2022 program funded by the NextGenerationEU (project ADVANCE - ADhesiVe free Fibre Metal LAminates fabricationN for aerospaCE applications - 2022W9SHCJ).
Before # 2, in order to overcome the higher friction force between the FML and dies due to the higher blank-holding force, a higher forming force is required. While the forming forces under Hf₁ and Hf₂ are almost the same after # 2, meaning that there is almost no relative sliding during the flattening process of the Z5 zone.

<figure><img src="image_473.png" /><figcaption>Fig. 6 Forming forces under different holding forces</figcaption></figure>

<figure><img src="image_474.png" /><figcaption>Fig. 7 Thickness distribution at # 1(a) 2 (b) and 3 (c) under the lowest blank-holding force.</figcaption></figure>

**Summary**

In this study, the thermoforming process was conducted on thermoplastic polymer and magnesium alloy-based FMLs to manufacture hat-shaped components. The effect of the blank-holding force on the forming results was analysed by comparing the forming force during the tests and the thickness distribution of each layer of the FML. The key findings can be summarized as follows:

* A non-uniform thickness distribution was found at different areas of the hat-shaped FML parts, and a significant reduction of the prepreg thickness was observed at the bottom radii compared to other areas.

* Moreover, the higher the blank-holding force the higher the forming force and the more significant the prepreg thickness variation.

**Acknowledgments**

This research was supported by Ministry University Research within the PRIN 2022 program funded by the NextGenerationEU (project ADVANCE - ADhesiVe free Fibre Metal LAminates fabricationN for aerospaCE applications - 2022W9SHCJ).
References

[1] X. Zhang, Q. Ma, Y. Dai, F. Hu, G. Liu, Z. Xu, G. Wei, T. Xu, Q. Zeng, W. Xie, Effects of surface treatments and bonding types on the interfacial behavior of fiber metal laminate based on magnesium alloy, Appl Surf Sci. 427 (2018) 897-906.
https://doi.org/10.1016/j.apsusc.2017.09.024

[2] Li B, Gong Y, Gao Y, Hong M, Li L. Failure analysis of hat-stringer-stiffened aircraft composite panels under fourpoint bending loading. Materials 2022;15(7):2430.

[3] Werner HO, Dörr D, Henning F, Kärger L. Numerical modeling of a hybrid forming process for three-dimensionally curved fiber-metal laminates. AIP Conference Proceedings 2019;2113:020019.

[4] Poppe CT, Werner HO, Kruse M, Chen H, Khalifa NB, Henning F, Kärger L. Towards 3D Process Simulation for In Situ Hybridization of Fiber-Metal-Laminates (FML), Key Engineering Materials 2022; 926:1399-1412.

[5] R. Alderliesten, On the development of hybrid material concepts for aircraft structures, Recent Pat. Eng. 3 (2009) 25-38. https://doi.org/10.2174/187221209787259893

[6] T. Sinmazcelik, E. Avcu, M.O. Bora, O. Coban, A review: fibre metal laminates, background, bonding types and applied test methods, Mater. Des. 32 (7) (2011) 3671-3686.
https://doi.org/10.1016/j.matdes.2011.03.011

[7] Y. Yang, R. Boom, R. Irion, D.-J. v. Heerden, P. Kuiper, H.d. Wit, Recycling of composite materials, Chem. Eng. Process 51 (2012) 53-68. https://doi.org/10.1016/j.cep.2011.09.007

[8] H. Chen, S. Li, J. Wang, A. Ding, A focused review on the thermo-stamping process and simulation progresses of continuous fibre reinforced thermoplastic composites, Compos B Eng. 224 (2021) 109196. https://doi.org/10.1016/j.compositesb.2021.109196

[9] Z Ding, H Wang, J Luo, N Li, A review on forming technologies of fibre metal laminates, International Journal of Lightweight Materials and Manufacture. 4 (2021) 110-126.
https://doi.org/10.1016/j.ijlmm.2020.06.006

[10] J.P.-H. Belnoue, O.J. Nixon-Pearson, D. Ivanov, S.R. Hallett, A novel hyper-viscoelastic model for consolidation of toughened prepregs under processing conditions, Mech Mater. 97 (2016) 118-134. https://doi.org/10.1016/j.mechmat.2016.02.019

[11] J.P.-H. Belnoue, M.A. Valverde, M. Onoufriou, X. Sun, D.S. Ivanov, S.R. Hallett, On the physical relevance of power law-based equations to describe the compaction behaviour of resin infused fibrous materials, Int J Mech Sci, 199 (2021) 106425.
https://doi.org/10.1016/j.ijmecsci.2021.106425

[12] A. Ghiotti, S. Bruschi, M. Kain, L. Lizzul, E. Simonetto, G. Tosellob, Simultaneous bonding and forming of Mg fibre metal laminates at high temperature. J Manuf Process; 72:105-114.
https://doi.org/10.1016/j.jmapro.2021.10.017

[13] T. Heggemann, W. Homberg, Deep drawing of fibre metal laminates for automotive lightweight structures, Compos. Struct. 216 (2019) 53-57.
https://doi.org/10.1016/j.compstruct.2019.02.047
## References

[1] X. Zhang, Q. Ma, Y. Dai, F. Hu, G. Liu, Z. Xu, G. Wei, T. Xu, Q. Zeng, W. Xie, Effects of surface treatments and bonding types on the interfacial behavior of fiber metal laminate based on magnesium alloy, Appl Surf Sci. 427 (2018) 897-906.
https://doi.org/10.1016/j.apsusc.2017.09.024

[2] Li B, Gong Y, Gao Y, Hong M, Li L. Failure analysis of hat-stringer-stiffened aircraft composite panels under fourpoint bending loading. Materials 2022;15(7):2430.

[3] Werner HO, Dörr D, Henning F, Kärger L. Numerical modeling of a hybrid forming process for three-dimensionally curved fiber-metal laminates. AIP Conference Proceedings 2019;2113:020019.

[4] Poppe CT, Werner HO, Kruse M, Chen H, Khalifa NB, Henning F, Kärger L. Towards 3D Process Simulation for In Situ Hybridization of Fiber-Metal-Laminates (FML), Key Engineering Materials 2022; 926:1399-1412.

[5] R. Alderliesten, On the development of hybrid material concepts for aircraft structures, Recent Pat. Eng. 3 (2009) 25-38. https://doi.org/10.2174/187221209787259893

[6] T. Sinmazcelik, E. Avcu, M.O. Bora, O. Coban, A review: fibre metal laminates, background, bonding types and applied test methods, Mater. Des. 32 (7) (2011) 3671-3686.
https://doi.org/10.1016/j.matdes.2011.03.011

[7] Y. Yang, R. Boom, R. Irion, D.-J. v. Heerden, P. Kuiper, H.d. Wit, Recycling of composite materials, Chem. Eng. Process 51 (2012) 53-68. https://doi.org/10.1016/j.cep.2011.09.007

[8] H. Chen, S. Li, J. Wang, A. Ding, A focused review on the thermo-stamping process and simulation progresses of continuous fibre reinforced thermoplastic composites, Compos B Eng. 224 (2021) 109196. https://doi.org/10.1016/j.compositesb.2021.109196

[9] Z Ding, H Wang, J Luo, N Li, A review on forming technologies of fibre metal laminates, International Journal of Lightweight Materials and Manufacture. 4 (2021) 110-126.
https://doi.org/10.1016/j.ijlmm.2020.06.006

[10] J.P.-H. Belnoue, O.J. Nixon-Pearson, D. Ivanov, S.R. Hallett, A novel hyper-viscoelastic model for consolidation of toughened prepregs under processing conditions, Mech Mater. 97 (2016) 118-134. https://doi.org/10.1016/j.mechmat.2016.02.019

[11] J.P.-H. Belnoue, M.A. Valverde, M. Onoufriou, X. Sun, D.S. Ivanov, S.R. Hallett, On the physical relevance of power law-based equations to describe the compaction behaviour of resin infused fibrous materials, Int J Mech Sci, 199 (2021) 106425.
https://doi.org/10.1016/j.ijmecsci.2021.106425

[12] A. Ghiotti, S. Bruschi, M. Kain, L. Lizzul, E. Simonetto, G. Tosellob, Simultaneous bonding and forming of Mg fibre metal laminates at high temperature. J Manuf Process; 72:105-114.
https://doi.org/10.1016/j.jmapro.2021.10.017

[13] T. Heggemann, W. Homberg, Deep drawing of fibre metal laminates for automotive lightweight structures, Compos. Struct. 216 (2019) 53-57.
https://doi.org/10.1016/j.compstruct.2019.02.047
# Influence of viscosity, binder activation, and loading rate on the membrane response of an infiltrated UD-NCF

MIRANDA PORTELA Renan¹,ª*, SCHÄFER Bastian²,ᵇ, KÄRGER Luise²,ᶜ, ROCHA DE FARIA Alfredo³,ᵈ, and MONTESANO John¹,ᵉ

¹Composites Research Group, Department of Mechanical & Mechatronics Engineering, University of Waterloo, Waterloo, Canada

²Karlsruhe Institute of Technology (KIT), Institute of Vehicle System Technology (FAST), Karlsruhe, Germany

³Department of Mechanical Engineering, Instituto Tecnológico de Aeronáutica (ITA), São José dos Campos, Brazil

ªrenan.portela@uwaterloo.ca, ᵇbastian.schaefer@kit.edu, ᶜluise.kaerger@kit.edu,

ᵈarfaria@ita.br, ᶜjohn.montesano@uwaterloo.ca

**Keywords:** Infiltrated Bias Extension Tests, Experimental Characterization, UD-NCF

**Abstract.** Shear-tension coupling of engineering fabric is one of the most important behaviors during the draping phase of liquid composite molding (LCM) processes, including wet compression molding (WCM), which occurs with the infiltrated fabric and, in some cases, with the use of a stabilizing binder. In the present study, the membrane behavior of an impregnated and binder-stabilized uni-directional carbon fiber non-crimp fabric was characterized by performing off-axis tension tests. These tests allow the investigation of the influence of stabilizing binder, fluid viscosity and loading rate on the fabric membrane behavior. As result of these experimental tests, an increase in membrane force is noticed when the stabilizing binder is activated, attributed to a greater shear stiffness. Additionally, a decrease in forces is observed for impregnated fabric compared to dry textiles caused by a lubrication layer between fiber tows. The study provides a better understanding of the membrane behavior of the impregnated and binder-stabilized UD-NCF, which is relevant for a potential high-volume production process.

**Introduction**

Automated wet compression molding (WCM) with fast-curing resin enables high-capacity production of continuous fiber-reinforced composite components [1]. The WCM process involves four simultaneously occurring physical mechanisms, which includes viscous draping, fluid progression, non-isothermal curing, and strong fluid-structure interaction [2]. Viscous draping of the fabric layers during mold closure is an inherent characteristic of the WCM process, involving complex mechanisms such as bending and membrane deformation, layer compaction, and frictions between adjacent layers and between layer and tool [2], [3]. During engineering textile draping process, membrane deformation is critical, as it can directly influence the formation of defects and misalignment of fibers [4], [5]. For the WCM process, membrane deformation may be influenced by the varying resin viscosity within the mold and the deformation rate.

Unidirectional non-crimp fabrics (UD-NCF) have been increasingly used in structural applications and, in some cases, are preferred over more conventional woven fabrics. UD-NCFs comprise parallel-oriented reinforcing fiber tows and low density transversely oriented supporting fibers held together by stitching thread, providing stability during the handling and manufacturing process. Unlike woven fabrics, UD-NCFs have no tow interlacing, minimizing tow crimp and, as a result, increasing the in-plane mechanical properties of fabricated composite parts [5]–[9]. However, the elimination of tow interlacing leads to lower fabric handleability, which may affect

<figure><img src="image_464.png" /></figure>
# Influence of viscosity, binder activation, and loading rate on the membrane response of an infiltrated UD-NCF

MIRANDA PORTELA Renan¹,ª*, SCHÄFER Bastian²,ᵇ, KÄRGER Luise²,ᶜ, ROCHA DE FARIA Alfredo³,ᵈ, and MONTESANO John¹,ᵉ

¹Composites Research Group, Department of Mechanical & Mechatronics Engineering, University of Waterloo, Waterloo, Canada

²Karlsruhe Institute of Technology (KIT), Institute of Vehicle System Technology (FAST), Karlsruhe, Germany

³Department of Mechanical Engineering, Instituto Tecnológico de Aeronáutica (ITA), São José dos Campos, Brazil

ªrenan.portela@uwaterloo.ca, ᵇbastian.schaefer@kit.edu, ᶜluise.kaerger@kit.edu, ᵈarfaria@ita.br, ᶮjohn.montesano@uwaterloo.ca

**Keywords:** Infiltrated Bias Extension Tests, Experimental Characterization, UD-NCF

**Abstract.** Shear-tension coupling of engineering fabric is one of the most important behaviors during the draping phase of liquid composite molding (LCM) processes, including wet compression molding (WCM), which occurs with the infiltrated fabric and, in some cases, with the use of a stabilizing binder. In the present study, the membrane behavior of an impregnated and binder-stabilized uni-directional carbon fiber non-crimp fabric was characterized by performing off-axis tension tests. These tests allow the investigation of the influence of stabilizing binder, fluid viscosity and loading rate on the fabric membrane behavior. As result of these experimental tests, an increase in membrane force is noticed when the stabilizing binder is activated, attributed to a greater shear stiffness. Additionally, a decrease in forces is observed for impregnated fabric compared to dry textiles caused by a lubrication layer between fiber tows. The study provides a better understanding of the membrane behavior of the impregnated and binder-stabilized UD-NCF, which is relevant for a potential high-volume production process.

**Introduction**

Automated wet compression molding (WCM) with fast-curing resin enables high-capacity production of continuous fiber-reinforced composite components [1]. The WCM process involves four simultaneously occurring physical mechanisms, which includes viscous draping, fluid progression, non-isothermal curing, and strong fluid-structure interaction [2]. Viscous draping of the fabric layers during mold closure is an inherent characteristic of the WCM process, involving complex mechanisms such as bending and membrane deformation, layer compaction, and frictions between adjacent layers and between layer and tool [2], [3]. During engineering textile draping process, membrane deformation is critical, as it can directly influence the formation of defects and misalignment of fibers [4], [5]. For the WCM process, membrane deformation may be influenced by the varying resin viscosity within the mold and the deformation rate.

Unidirectional non-crimp fabrics (UD-NCF) have been increasingly used in structural applications and, in some cases, are preferred over more conventional woven fabrics. UD-NCFs comprise parallel-oriented reinforcing fiber tows and low density transversely oriented supporting fibers held together by stitching thread, providing stability during the handling and manufacturing process. Unlike woven fabrics, UD-NCFs have no tow interlacing, minimizing tow crimp and, as a result, increasing the in-plane mechanical properties of fabricated composite parts [5]–[9]. However, the elimination of tow interlacing leads to lower fabric handleability, which may affect

<figure><img src="image_475.png" /></figure>
the manufacturing automation and process time [10]. The addition of a stabilizing binder can improve the handling of UD-NCFs by providing higher cohesion between adjacent layers and mitigating ply slippage, while potentially reducing the manufacturing cost [5], [10]. For thick WCM parts with complex geometries, binder stabilized UD-NCFs may aid in increasing the fabric stack stiffness for improved resin management prior to mold closure. For binder stabilized UD-NCFs, the pre-activation of the binder may further influence the membrane deformation.

In general, membrane deformation is measured either by picture frame test (PFT) or bias-extension test (BET) setups [11], [12]. The PFT works by applying uniaxial tensile forces to the low and top corners of the testing frame. It is conducted with a testing frame, initially in a square shape. During the experimental test, the testing frame deforms into a rhomboid while simultaneously imposing pure shear stress on the fabric sample. PFTs have been used for characterizing the shear deformation of woven fabrics due to their balanced nature [4], [13]. However, PFTs are not suitable for the characterization of UD-NCFs due to the fabric behavior that combines pure and simple shear [14], [15]. These two shear mechanisms are represented by the rotation and slippage of the fiber tows, respectively [14]. On the other hand, BET is performed with a rectangular shape fabric specimen that must have length to width ratio greater than 2. This type of test can be used to measure the extension-shear behavior with the fabric oriented in different directions, such as 30°, 45°, and 60°. In Ref [16], the authors indicated that 45° BET can be used to capture the UD-NCF shear response, while other studies suggest that 30° and 60° BET provide a more comprehensive understanding of the UD-NCF behavior [12], [14], [15].

Both the PFT and BET are suitable to characterize the membrane deformation of dry fabric. In Refs. [1], [2], an infiltrated bias-extension test (IBET) setup was designed to characterize the extension-shear deformation of impregnated woven fabric, which is representative of fabric draping during the WCM process. The IBET requires a horizontal frame with a reservoir, a rear fixture that holds the fabric, and a moving bridge that extends the engineering textile. In the reservoir, an amount of silicone oil is poured to fully impregnate the fabric before the test. Although impregnated woven fabrics were characterized in previous work [1], [2], the extension-shear behavior of impregnated UD-NCFs has not been studied. Furthermore, the influence of binder activation on the membrane deformation of UD-NCFs has not been studied.

In this work, the shear-tension behavior of an impregnated and binder stabilized carbon fiber UD-NCF is investigated through a series of IBETs. Thereby, the influence of the loading rate, fluid viscosity, and activation of binder on the membrane deformation response is investigated.

## Experimental Procedure

*Materials.* A binder stabilized UD-NCF (Figure 1) was investigated, namely PX35-UD300 (Zoltek, US). This fabric consists of a single layer of axially oriented parallel tows each comprising 50K continuous PX35 carbon fibers (CFs) stitched together with a 76 dtex polyester yarn in a tricot pattern. The fabric also consists of transversely oriented low linear density supporting glass fibers (GF) for improved handleability and a thermosetting polymer powder binder applied on one surface. The total areal density of the UD-NCF is 333 g/m². To impregnate the textiles and avoid additional temperature-related uncertainties, silicone oils with two different viscosities (20 and 100 cSt) were used.
the manufacturing automation and process time [10]. The addition of a stabilizing binder can improve the handling of UD-NCFs by providing higher cohesion between adjacent layers and mitigating ply slippage, while potentially reducing the manufacturing cost [5], [10]. For thick WCM parts with complex geometries, binder stabilized UD-NCFs may aid in increasing the fabric stack stiffness for improved resin management prior to mold closure. For binder stabilized UD-NCFs, the pre-activation of the binder may further influence the membrane deformation.

In general, membrane deformation is measured either by picture frame test (PFT) or bias-extension test (BET) setups [11], [12]. The PFT works by applying uniaxial tensile forces to the low and top corners of the testing frame. It is conducted with a testing frame, initially in a square shape. During the experimental test, the testing frame deforms into a rhomboid while simultaneously imposing pure shear stress on the fabric sample. PFTs have been used for characterizing the shear deformation of woven fabrics due to their balanced nature [4], [13]. However, PFTs are not suitable for the characterization of UD-NCFs due to the fabric behavior that combines pure and simple shear [14], [15]. These two shear mechanisms are represented by the rotation and slippage of the fiber tows, respectively [14]. On the other hand, BET is performed with a rectangular shape fabric specimen that must have length to width ratio greater than 2. This type of test can be used to measure the extension-shear behavior with the fabric oriented in different directions, such as 30°, 45°, and 60°. In Ref [16], the authors indicated that 45° BET can be used to capture the UD-NCF shear response, while other studies suggest that 30° and 60° BET provide a more comprehensive understanding of the UD-NCF behavior [12], [14], [15].

Both the PFT and BET are suitable to characterize the membrane deformation of dry fabric. In Refs. [1], [2], an infiltrated bias-extension test (IBET) setup was designed to characterize the extension-shear deformation of impregnated woven fabric, which is representative of fabric draping during the WCM process. The IBET requires a horizontal frame with a reservoir, a rear fixture that holds the fabric, and a moving bridge that extends the engineering textile. In the reservoir, an amount of silicone oil is poured to fully impregnate the fabric before the test. Although impregnated woven fabrics were characterized in previous work [1], [2], the extension-shear behavior of impregnated UD-NCFs has not been studied. Furthermore, the influence of binder activation on the membrane deformation of UD-NCFs has not been studied.

In this work, the shear-tension behavior of an impregnated and binder stabilized carbon fiber UD-NCF is investigated through a series of IBETs. Thereby, the influence of the loading rate, fluid viscosity, and activation of binder on the membrane deformation response is investigated.

## Experimental Procedure

*Materials.* A binder stabilized UD-NCF (Figure 1) was investigated, namely PX35-UD300 (Zoltek, US). This fabric consists of a single layer of axially oriented parallel tows each comprising 50K continuous PX35 carbon fibers (CFs) stitched together with a 76 dtex polyester yarn in a tricot pattern. The fabric also consists of transversely oriented low linear density supporting glass fibers (GF) for improved handleability and a thermosetting polymer powder binder applied on one surface. The total areal density of the UD-NCF is 333 g/m². To impregnate the textiles and avoid additional temperature-related uncertainties, silicone oils with two different viscosities (20 and 100 cSt) were used.
<figure><img src="image_465.png" /><figcaption>Figure 1: Binder stabilized uni-directional non-crimp fabric (PX35-UD300); left: Polyester stitching yarns in a tricot pattern; right: Glass fiber to improve the fabric handleability.</figcaption></figure>

Infiltrated bias-extension test (IBET) The IBET was performed on a servo-hydraulic test frame within the same fixture used by Poppe et al. [1], comprising a reservoir, a moving bridge attached to a 1 mm wire and a rear fixture as illustrated in Figure 2. When the IBET is conducted, the moving bridge is pulled, and the specimen, attached to both fixtures, is stretched. This setup was mounted on a servo-hydraulic test frame with a 1 kN load cell. Undeformed layers with a size of 280 × 140 mm² and two different bias angles (30° and 60°) were tested at different loading rates of 100 and 500 mm/min, under both dry and infiltrated conditions with two different liquid viscosities (20 cSt and 100 cSt), and with the binder pre-activated and inactivated. Each test was performed three times with a new UD-NCF specimen.

<figure><img src="image_466.png" /><figcaption>Figure 2: Experimental IBET setup.</figcaption></figure>

Similar to the evaluation methodology proposed by Poppe et al. [1], the shear resulting forces ($f_{shear}$) were found by:

$$f_{\text{shear}} = f_{\text{raw}} - f_{\text{friction}} - f_{\text{fluid}}(v, \mu) \quad (1)$$

where $f_{raw}$ is the force measured by the universal testing machine, $f_{friction}$ is the frictional force within the linear bearings, and $f_{fluid}(v, \mu)$ is the fluid resistance on the moving, which is expressed as a function of fluid viscosity ($\mu$) and loading rate or velocity ($v$). The last two forces were
<figure><img src="image_476.png" /><figcaption>Figure 1: Binder stabilized uni-directional non-crimp fabric (PX35-UD300); left: Polyester stitching yarns in a tricot pattern; right: Glass fiber to improve the fabric handleability.</figcaption></figure>

Infiltrated bias-extension test (IBET) The IBET was performed on a servo-hydraulic test frame within the same fixture used by Poppe et al. [1], comprising a reservoir, a moving bridge attached to a 1 mm wire and a rear fixture as illustrated in Figure 2. When the IBET is conducted, the moving bridge is pulled, and the specimen, attached to both fixtures, is stretched. This setup was mounted on a servo-hydraulic test frame with a 1 kN load cell. Undeformed layers with a size of 280 × 140 mm² and two different bias angles (30° and 60°) were tested at different loading rates of 100 and 500 mm/min, under both dry and infiltrated conditions with two different liquid viscosities (20 cSt and 100 cSt), and with the binder pre-activated and inactivated. Each test was performed three times with a new UD-NCF specimen.

<figure><img src="image_477.png" /><figcaption>Figure 2: Experimental IBET setup.</figcaption></figure>

Similar to the evaluation methodology proposed by Poppe et al. [1], the shear resulting forces ($f_{shear}$) were found by:

$$f_{\text{shear}} = f_{\text{raw}} - f_{\text{friction}} - f_{\text{fluid}}(v, \mu) \quad (1)$$

where $f_{raw}$ is the force measured by the universal testing machine, $f_{friction}$ is the frictional force within the linear bearings, and $f_{fluid}(v, \mu)$ is the fluid resistance on the moving, which is expressed as a function of fluid viscosity ($\mu$) and loading rate or velocity ($v$). The last two forces were
determined by performing the IBET with no specimen. For each configuration of viscosity and velocity, calibration tests were conducted 5 times.

## Results and Discussion

*Influence of fluid viscosity.* To investigate the influence of the fluid viscosity, two different fabric orientations (30° and 60°) were tested, both dry and impregnated with two different fluid viscosities (20 and 100 cSt). The displacement fields of both fabric orientations are illustrated in Figure 3. The results for the 30°-IBET and 60°-IBET are shown in Figure 4 and 5, respectively. For both considered fluid viscosities in the 30°-IBET, the average force of the dry specimens was up to 30% higher at the final displacement of approximately 50 mm (Figure 4). When the fabric is impregnated, it is observed a membrane force reduction, which is attributed to a lower friction between fibers within the tow and between the fiber tows and the stitching, resulting from the lubrication layer on the fibers tows [1]. The same phenomenon occurs for an initial orientation of 60°, cf. Figure 5 (right), where the dry specimen has a higher membrane force. This force is reduced when the specimen is impregnated. On the other hand, an anomaly was observed in Figure 5 (left), showing higher membrane forces in the infiltrated specimens with 100 cSt compared to the dry specimens. This may be due to slippage of the sample during the experiment.

<figure><img src="image_467.png" /><figcaption>Figure 3: UD-NCF deformation field; left: 30° test specimen; right: 60° test specimen.</figcaption></figure>
determined by performing the IBET with no specimen. For each configuration of viscosity and velocity, calibration tests were conducted 5 times.

## Results and Discussion

*Influence of fluid viscosity.* To investigate the influence of the fluid viscosity, two different fabric orientations (30° and 60°) were tested, both dry and impregnated with two different fluid viscosities (20 and 100 cSt). The displacement fields of both fabric orientations are illustrated in Figure 3. The results for the 30°-IBET and 60°-IBET are shown in Figure 4 and 5, respectively. For both considered fluid viscosities in the 30°-IBET, the average force of the dry specimens was up to 30% higher at the final displacement of approximately 50 mm (Figure 4). When the fabric is impregnated, it is observed a membrane force reduction, which is attributed to a lower friction between fibers within the tow and between the fiber tows and the stitching, resulting from the lubrication layer on the fibers tows [1]. The same phenomenon occurs for an initial orientation of 60°, cf. Figure 5 (right), where the dry specimen has a higher membrane force. This force is reduced when the specimen is impregnated. On the other hand, an anomaly was observed in Figure 5 (left), showing higher membrane forces in the infiltrated specimens with 100 cSt compared to the dry specimens. This may be due to slippage of the sample during the experiment.

<figure><img src="image_478.png" /><figcaption>Figure 3: UD-NCF deformation field; left: 30° test specimen; right: 60° test specimen.</figcaption></figure>
<figure><img src="image_468.png" /><figcaption>Figure 4: Experimental results for dry (blue) and impregnate (20 cSt in orange and 100 cSt in green) UD-NCF specimens; left: 30° specimens tested at 100 mm/min; right: 30° specimens tested at 500 mm/min.</figcaption></figure>

<figure><img src="image_469.png" /><figcaption>Figure 5: Experimental results for dry (blue) and impregnated (20 cSt in orange and 100 cSt in green) UD-NCF specimens; left: 60° specimens tested at 100 mm/min; right: 60° specimens tested at 500 mm/min.</figcaption></figure>

Influence of fabric binder pre-activation. Figure 6 presents the fabric with inactivated binder (left) and the pre-activated binder (right). Figure 6 (left) shows the presence of polymer grains, which are pulverized over the fabric surface by the material supplier during the textile manufacturing process. When the binder is pre-activated in the oven at 120°C, these grains melt and coat the nearest fibers and stitching. Figures 7 and 8 illustrate the influence of the pre-activated binder on the membrane force of the fabric initially oriented at 30° and 60°, respectively. The configurations with pre-activated binder are denoted by AT and without pre-activation of the binder as IT. The membrane force increases as the stabilizing binder is activated. This is due to the polymer binder
<figure><img src="image_479.png" /><figcaption>Figure 4: Experimental results for dry (blue) and impregnate (20 cSt in orange and 100 cSt in green) UD-NCF specimens; left: 30° specimens tested at 100 mm/min; right: 30° specimens tested at 500 mm/min.</figcaption></figure>

<figure><img src="image_480.png" /><figcaption>Figure 5: Experimental results for dry (blue) and impregnated (20 cSt in orange and 100 cSt in green) UD-NCF specimens; left: 60° specimens tested at 100 mm/min; right: 60° specimens tested at 500 mm/min.</figcaption></figure>

Influence of fabric binder pre-activation. Figure 6 presents the fabric with inactivated binder (left) and the pre-activated binder (right). Figure 6 (left) shows the presence of polymer grains, which are pulverized over the fabric surface by the material supplier during the textile manufacturing process. When the binder is pre-activated in the oven at 120°C, these grains melt and coat the nearest fibers and stitching. Figures 7 and 8 illustrate the influence of the pre-activated binder on the membrane force of the fabric initially oriented at 30° and 60°, respectively. The configurations with pre-activated binder are denoted by AT and without pre-activation of the binder as IT. The membrane force increases as the stabilizing binder is activated. This is due to the polymer binder
that melts and coats the fiber tows and additionally increases cohesion between the tows and the stitching. This increases the shear stiffness of the fabric. In some cases, this phenomenon increases the membrane force up to 100% at different displacement values, including 20 mm and 40 mm, as shown in Figure 8 (right).

<figure><img src="image_470.png" /><figcaption>Figure 6: Binder-stabilized UD-NCF; left: Inactivated binder; right: Pre-activated binder in the oven at 120°C.</figcaption></figure>

<figure><img src="image_471.png" /><figcaption>Figure 7: Experimental results for dry inactivated (IT) and pre-activated (AT) binder on the UD-NCF specimens; left: 30° specimens tested at 100 mm/min; right: 30° specimens tested at 500 mm/min.</figcaption></figure>
that melts and coats the fiber tows and additionally increases cohesion between the tows and the stitching. This increases the shear stiffness of the fabric. In some cases, this phenomenon increases the membrane force up to 100% at different displacement values, including 20 mm and 40 mm, as shown in Figure 8 (right).

<figure><img src="image_481.png" /><figcaption>Figure 6: Binder-stabilized UD-NCF; left: Inactivated binder; right: Pre-activated binder in the oven at 120°C.</figcaption></figure>

<figure><img src="image_482.png" /><figcaption>Figure 7: Experimental results for dry inactivated (IT) and pre-activated (AT) binder on the UD-NCF specimens; left: 30° specimens tested at 100 mm/min; right: 30° specimens tested at 500 mm/min.</figcaption></figure>
<figure><img src="image_472.png" /><figcaption>Figure 8: Experimental results for dry inactivated (IT) and pre-activated (AT) binder on the UD-NCF specimens; left: 60° specimens tested at 100 mm/min; right: 60° specimens tested at 500 mm/min.</figcaption></figure>

**Influence of loading rate.** To understand the influence of the loading rate, the specimens were tested at two different rates. The 30°- and 60°-IBET specimens, presented in Figures 9 and 10, respectively, were tested under 100 and 500 $\frac{\text{mm}}{\text{min}}$. Figure 9 illustrates the results of the 30° specimens that overlap regardless of the infiltration state of the fabric, which means that for this orientation the loading rate has no influence over the membrane force. The results of the 60°-IBET specimens show an increase of the average membrane force when the specimens were dry (Figure 10 left); however, the standard deviation of both loading rates overlap, representing no significant difference between both loading rates. Analogously, the results of the infiltrated samples with 20 cSt silicone oil are similar, indicating no influence of the loading rate on the membrane force. Lastly, a higher membrane force was obtained at a slower loading rate when the specimens were infiltrated with 100 cSt silicone oil. The 60° bias angle is more complex because of the possible slippage of the specimen which can have affected the results presented in Figure 10 right.

<figure><img src="image_473.png" /><figcaption>Figure 9: Experimental results for dry and impregnated UD-NCF specimens with activated binder system (AT); left: dry 30° specimens tested at 100 and 500 mm/min; center: impregnated 30° specimens under 20 cSt tested at 100 and 500 mm/min; right: impregnated 30° specimens under 100 cSt tested at 100 and 500 mm/min.</figcaption></figure>
<figure><img src="image_483.png" /><figcaption>Figure 8: Experimental results for dry inactivated (IT) and pre-activated (AT) binder on the UD-NCF specimens; left: 60° specimens tested at 100 mm/min; right: 60° specimens tested at 500 mm/min.</figcaption></figure>

**Influence of loading rate.** To understand the influence of the loading rate, the specimens were tested at two different rates. The 30°- and 60°-IBET specimens, presented in Figures 9 and 10, respectively, were tested under 100 and 500 $\frac{\text{mm}}{\text{min}}$. Figure 9 illustrates the results of the 30° specimens that overlap regardless of the infiltration state of the fabric, which means that for this orientation the loading rate has no influence over the membrane force. The results of the 60°-IBET specimens show an increase of the average membrane force when the specimens were dry (Figure 10 left); however, the standard deviation of both loading rates overlap, representing no significant difference between both loading rates. Analogously, the results of the infiltrated samples with 20 cSt silicone oil are similar, indicating no influence of the loading rate on the membrane force. Lastly, a higher membrane force was obtained at a slower loading rate when the specimens were infiltrated with 100 cSt silicone oil. The 60° bias angle is more complex because of the possible slippage of the specimen which can have affected the results presented in Figure 10 right.

<figure><img src="image_484.png" /><figcaption>Figure 9: Experimental results for dry and impregnated UD-NCF specimens with activated binder system (AT); left: dry 30° specimens tested at 100 and 500 mm/min; center: impregnated 30° specimens under 20 cSt tested at 100 and 500 mm/min; right: impregnated 30° specimens under 100 cSt tested at 100 and 500 mm/min.</figcaption></figure>
<figure><img src="image_474.png" /><figcaption>Figure 10: Experimental results for dry and impregnated UD-NCF specimens with activated binder system (AT); left: dry 60° specimens tested at 100 and 500 mm/min; center: impregnated 60° specimens under 20 cSt tested at 100 and 500 mm/min; right: impregnated 60° specimens under 100 cSt tested at 100 and 500 mm/min.</figcaption></figure>

## Summary and conclusion

In conclusion, the binder activation causes an increase of the fabric's membrane stiffness. The increase of the membrane stiffness happens because of the polymer binder that melts and, consequently, coats the fiber tows and increases cohesion with the stitching, increasing the fabric's membrane stiffness. On the other hand, the fluid viscosity creates a lubrication layer between the fiber tows and the stitching, reducing the friction, and consequently, decreasing the membrane stiffness. At last, no significant influence of the strain rate on the 30°- and 60°-IBET specimen results was found.

## References

[1] C. Poppe, D. Dörr, F. Henning, and L. Kärger, "Experimental and numerical investigation of the shear behaviour of infiltrated woven fabrics," *Compos. Part A Appl. Sci. Manuf.*, vol. 114, pp. 327-337, Nov. 2018. https://doi.org/10.1016/j.compositesa.2018.08.018

[2] C. Poppe, "Process simulation of wet compression moulding for continuous fibre-reinforced polymers," Karlsruhe Institute of Technology, Karlsruhe, 2021.

[3] C. T. Poppe, C. Krauß, F. Albrecht, and L. Kärger, "A 3D process simulation model for wet compression moulding," *Compos. Part A Appl. Sci. Manuf.*, vol. 145, p. 106379, 2021. https://doi.org/10.1016/j.compositesa.2021.106379

[4] M. Ghazimoradi, E. A. Trejo, C. Butcher, and J. Montesano, "Characterizing the macroscopic response and local deformation mechanisms of a unidirectional non-crimp fabric," *Compos. Part A Appl. Sci. Manuf.*, vol. 156, p. 106857, May 2022. https://doi.org/10.1016/j.compositesa.2022.106857

[5] J. Pourtier, B. Duchamp, M. Kowalski, X. Legrand, P. Wang, and D. Soulat, "Bias extension test on a bi-axial non-crimp fabric powdered with a non-reactive binder system," in *AIP Conference Proceedings*, 2018, vol. 1960, p. 020023. https://doi.org/10.1063/1.5034824

[6] S. V. Lomov, *Non-crimp fabric composites*. Cambridge, UK: Woodhead Publishing Limited, 2011. https://doi.org/10.1533/9780857092533

[7] B. Schäfer, R. Zheng, P. Boisse, and L. Kärger, "Investigation of the compaction behavior of uni-and bidirectional non-crimp fabrics," in *Materials Research Proceedings*, 2023, vol. 28, pp. 331-338. https://doi.org/10.21741/9781644902479-36
<figure><img src="image_485.png" /><figcaption>Figure 10: Experimental results for dry and impregnated UD-NCF specimens with activated binder system (AT); left: dry 60° specimens tested at 100 and 500 mm/min; center: impregnated 60° specimens under 20 cSt tested at 100 and 500 mm/min; right: impregnated 60° specimens under 100 cSt tested at 100 and 500 mm/min.</figcaption></figure>

## Summary and conclusion

In conclusion, the binder activation causes an increase of the fabric's membrane stiffness. The increase of the membrane stiffness happens because of the polymer binder that melts and, consequently, coats the fiber tows and increases cohesion with the stitching, increasing the fabric's membrane stiffness. On the other hand, the fluid viscosity creates a lubrication layer between the fiber tows and the stitching, reducing the friction, and consequently, decreasing the membrane stiffness. At last, no significant influence of the strain rate on the 30°- and 60°-IBET specimen results was found.

## References

[1] C. Poppe, D. Dörr, F. Henning, and L. Kärger, “Experimental and numerical investigation of the shear behaviour of infiltrated woven fabrics,” *Compos. Part A Appl. Sci. Manuf.*, vol. 114, pp. 327–337, Nov. 2018. https://doi.org/10.1016/j.compositesa.2018.08.018

[2] C. Poppe, “Process simulation of wet compression moulding for continuous fibre-reinforced polymers,” Karlsruhe Institute of Technology, Karlsruhe, 2021.

[3] C. T. Poppe, C. Krauß, F. Albrecht, and L. Kärger, “A 3D process simulation model for wet compression moulding,” *Compos. Part A Appl. Sci. Manuf.*, vol. 145, p. 106379, 2021. https://doi.org/10.1016/j.compositesa.2021.106379

[4] M. Ghazimoradi, E. A. Trejo, C. Butcher, and J. Montesano, “Characterizing the macroscopic response and local deformation mechanisms of a unidirectional non-crimp fabric,” *Compos. Part A Appl. Sci. Manuf.*, vol. 156, p. 106857, May 2022. https://doi.org/10.1016/j.compositesa.2022.106857

[5] J. Pourtier, B. Duchamp, M. Kowalski, X. Legrand, P. Wang, and D. Soulat, “Bias extension test on a bi-axial non-crimp fabric powdered with a non-reactive binder system,” in *AIP Conference Proceedings*, 2018, vol. 1960, p. 020023. https://doi.org/10.1063/1.5034824

[6] S. V. Lomov, *Non-crimp fabric composites*. Cambridge, UK: Woodhead Publishing Limited, 2011. https://doi.org/10.1533/9780857092533

[7] B. Schäfer, R. Zheng, P. Boisse, and L. Kärger, “Investigation of the compaction behavior of uni-and bidirectional non-crimp fabrics,” in *Materials Research Proceedings*, 2023, vol. 28, pp. 331–338. https://doi.org/10.21741/9781644902479-36
[8] P. Quenzel et al., "Material characterisation of biaxial glass-fibre non-crimp fabrics as a function of ply orientation, stitch pattern, stitch length and stitch tension," *J. Compos. Mater.*, vol. 56, no. 26, pp. 3971-3991, Nov. 2022. https://doi.org/10.1177/00219983221127005

[9] L. Kärger, S. Galkin, E. Kunze, M. Gude, and B. Schäfer, "Prediction of forming effects in UD-NCF by macroscopic forming simulation – Capabilities and limitations," in *ESAFORM 2021*, 2021, pp. 14–16. https://doi.org/10.25518/esaform21.355

[10] P. H. Broberg, C. Krogh, E. Lindgaard, and B. L. V. Bak, "Simulation of Wrinkling during Forming of Binder Stabilized UD-NCF Preforms in Wind Turbine Blade Manufacturing," *Key Eng. Mater.*, vol. 926, pp. 1248–1256, Jul. 2022. https://doi.org/10.4028/p-165q46

[11] P. Boisse, N. Hamila, E. Guzman-Maldonado, A. Madeo, G. Hivet, and F. Dell'Isola, "The bias-extension test for the analysis of in-plane shear properties of textile composite reinforcements and prepregs: a review," *Int. J. Mater. Form.*, vol. 10, no. 4, pp. 473–492, Aug. 2017. https://doi.org/10.1007/s12289-016-1294-7

[12] F. J. Schirmaier, D. Dörr, F. Henning, and L. Kärger, "A macroscopic approach to simulate the forming behaviour of stitched unidirectional non-crimp fabrics (UD-NCF)," *Compos. Part A Appl. Sci. Manuf.*, vol. 102, pp. 322–335, Nov. 2017. https://doi.org/10.1016/j.compositesa.2017.08.009

[13] G. Barbagallo, A. Madeo, I. Azehaf, I. Giorgio, F. Morestin, and P. Boisse, "Bias extension test on an unbalanced woven composite reinforcement: Experiments and modeling via a second-gradient continuum approach," *J. Compos. Mater.*, vol. 51, no. 2, pp. 153–170, Jan. 2017. https://doi.org/10.1177/0021998316643577

[14] B. Schäfer, R. Zheng, N. Naouar, and L. Kärger, "Membrane behavior of uni- and bidirectional non-crimp fabrics in off-axis-tension tests," *Int. J. Mater. Form.*, vol. 16, no. 6, pp. 16–20, 2023. https://doi.org/10.1007/s12289-023-01792-x

[15] F. J. Schirmaier, K. A. Weidenmann, L. Kärger, and F. Henning, "Characterisation of the draping behaviour of unidirectional non-crimp fabrics (UD-NCF)," *Compos. Part A Appl. Sci. Manuf.*, vol. 80, pp. 28–38, Jan. 2016. https://doi.org/10.1016/j.compositesa.2015.10.004

[16] M. Ghazimoradi and J. Montesano, "Macroscopic Forming Simulation for a Unidirectional Non-crimp Fabric Using an Anisotropic Hyperelastic Material Model," *Appl. Compos. Mater.*, vol. 30, no. 6, pp. 2001–2023, Dec. 2023. https://doi.org/10.1007/s10443-023-10158-0
[8] P. Quenzel et al., "Material characterisation of biaxial glass-fibre non-crimp fabrics as a function of ply orientation, stitch pattern, stitch length and stitch tension," *J. Compos. Mater.*, vol. 56, no. 26, pp. 3971-3991, Nov. 2022. https://doi.org/10.1177/00219983221127005

[9] L. Kärger, S. Galkin, E. Kunze, M. Gude, and B. Schäfer, "Prediction of forming effects in UD-NCF by macroscopic forming simulation – Capabilities and limitations," in *ESAFORM 2021*, 2021, pp. 14–16. https://doi.org/10.25518/esaform21.355

[10] P. H. Broberg, C. Krogh, E. Lindgaard, and B. L. V. Bak, "Simulation of Wrinkling during Forming of Binder Stabilized UD-NCF Preforms in Wind Turbine Blade Manufacturing," *Key Eng. Mater.*, vol. 926, pp. 1248–1256, Jul. 2022. https://doi.org/10.4028/p-165q46

[11] P. Boisse, N. Hamila, E. Guzman-Maldonado, A. Madeo, G. Hivet, and F. Dell'Isola, "The bias-extension test for the analysis of in-plane shear properties of textile composite reinforcements and prepregs: a review," *Int. J. Mater. Form.*, vol. 10, no. 4, pp. 473–492, Aug. 2017. https://doi.org/10.1007/s12289-016-1294-7

[12] F. J. Schirmaier, D. Dörr, F. Henning, and L. Kärger, "A macroscopic approach to simulate the forming behaviour of stitched unidirectional non-crimp fabrics (UD-NCF)," *Compos. Part A Appl. Sci. Manuf.*, vol. 102, pp. 322–335, Nov. 2017. https://doi.org/10.1016/j.compositesa.2017.08.009

[13] G. Barbagallo, A. Madeo, I. Azehaf, I. Giorgio, F. Morestin, and P. Boisse, "Bias extension test on an unbalanced woven composite reinforcement: Experiments and modeling via a second-gradient continuum approach," *J. Compos. Mater.*, vol. 51, no. 2, pp. 153–170, Jan. 2017. https://doi.org/10.1177/0021998316643577

[14] B. Schäfer, R. Zheng, N. Naouar, and L. Kärger, "Membrane behavior of uni- and bidirectional non-crimp fabrics in off-axis-tension tests," *Int. J. Mater. Form.*, vol. 16, no. 6, pp. 16–20, 2023. https://doi.org/10.1007/s12289-023-01792-x

[15] F. J. Schirmaier, K. A. Weidenmann, L. Kärger, and F. Henning, "Characterisation of the draping behaviour of unidirectional non-crimp fabrics (UD-NCF)," *Compos. Part A Appl. Sci. Manuf.*, vol. 80, pp. 28–38, Jan. 2016. https://doi.org/10.1016/j.compositesa.2015.10.004

[16] M. Ghazimoradi and J. Montesano, "Macroscopic Forming Simulation for a Unidirectional Non-crimp Fabric Using an Anisotropic Hyperelastic Material Model," *Appl. Compos. Mater.*, vol. 30, no. 6, pp. 2001–2023, Dec. 2023. https://doi.org/10.1007/s10443-023-10158-0
# Polymer-metal interactions and their effect on tool-ply friction of C/PEKK in melt

GROUVE Wouter<sup>1,a*</sup>, PAMUNGKAS Hafidz<sup>1,b</sup>, HELTHUIS Nick<sup>1,c</sup>, PIERIK Rens<sup>1,2,d</sup>, AKKERMAN Remko<sup>1,2,e</sup>

<sup>1</sup>Chair of Production Technology, Faculty of Engineering Technology, University of Twente, Drienerlolaan 5, 7622 NB Enschede, the Netherlands

<sup>2</sup>ThermoPlastic composites Research Center, Palatijn 15, 7521 PN Enschede, the Netherlands

<sup>a</sup>w.j.b.grouve@utwente.nl, <sup>b</sup>hafidzelamiennoorpamungkas@student.utwente.nl,

<sup>c</sup>n.g.j.helthuis@utwente.nl, <sup>d</sup>e.r.pierik@utwente.nl, <sup>e</sup>r.akkerman@uwtwente.nl

**Keywords:** Thermoplastic Matrix, Tool-Ply Friction, Chemical Bonding, Polymer Degradation

**Abstract.** Excessive friction or adhesion between tool and surface ply can lead to processing defects during thermoplastic composite manufacturing. A previous study identified a notable increase in friction or adhesion between a mild steel tool surface and C/PEKK with prolonged dwell time at elevated temperature, a phenomenon that could not be explained by our current understanding. This work presents additional friction and differential scanning calorimetry experiments to demonstrate that the increased adhesion results from polymer-metal interactions that lead to polymer degradation and metal-polymer bonding. These findings underline the importance of tool material selection in both processing and characterization of tool-ply friction.

**Introduction**

High friction or adhesion between tool surface and composite material can lead to fiber waviness defects during the manufacture of thermoplastic composite parts [1-3]. Many research groups have characterized tool-ply friction under processing conditions in an effort to understand and prevent such defect formation. The characterization experiments involve the controlled sliding of a molten thermoplastic composite ply against a representative tool surface [4]. The measured shear stress at the interface typically displays a peak, followed by a steady-state plateau, where the peak and steady-state stress levels depend on temperature, applied normal pressure and sliding rate [5-7].

The mechanisms underlying the observed shear stress response are schematically illustrated in Fig. 1, and include friction, physical and/or chemical bonding, and viscoelastic deformation of the molten polymer. The peak shear stress could be caused by an adhesive bond or static friction between the fibers and the tool, which needs to be overcome before sliding can commence [8, 9]. Additionally, a thin polymer interlayer exists between the fibers and the metal surface, acting as a lubricating film. Here, the evolution of the shear stress is governed by the deformation of the matrix interlayer. The peak stress can then be caused by the nonlinear viscoelasticity of the matrix material [6, 7], a change in the interlayer thickness distribution [5], or by the development of slip at the tool-polymer [5] and/or the fiber-polymer interface [10, 11].

Although the current understanding of these phenomena is well-developed, our previous research [12] brought to light a curious phenomenon: the peak friction between mild steel and Solvay unidirectional C/PEKK significantly increased with dwell time at elevated temperature; an effect that was not observed for ply-ply friction. This unexpected finding challenged our understanding of the governing mechanisms and prompted additional experiments to identify the underlying causes. In this paper, we report results indicating that the increased peak friction is caused by strong polymer-steel bonding and polymer cross-linking, with the latter accelerated by

<figure><img src="image_475.png" /></figure>
# Polymer-metal interactions and their effect on tool-ply friction of C/PEKK in melt

GROUVE Wouter¹,ᵃ*, PAMUNGKAS Hafidz¹,ᵇ, HELTHUIS Nick¹,ᶜ, PIERIK Rens¹,²,ᵈ, AKKERMAN Remko¹,²,ᵉ

¹Chair of Production Technology, Faculty of Engineering Technology, University of Twente, Drienerlolaan 5, 7622 NB Enschede, the Netherlands

²ThermoPlastic composites Research Center, Palatijn 15, 7521 PN Enschede, the Netherlands

ᵃw.j.b.grouve@utwente.nl, ᵇhafidzelamiennoorpamungkas@student.utwente.nl,

ᶜn.g.j.helthuis@utwente.nl, ᵈe.r.pierik@utwente.nl, ᶜr.akkerman@uwtwente.nl

**Keywords:** Thermoplastic Matrix, Tool-Ply Friction, Chemical Bonding, Polymer Degradation

**Abstract.** Excessive friction or adhesion between tool and surface ply can lead to processing defects during thermoplastic composite manufacturing. A previous study identified a notable increase in friction or adhesion between a mild steel tool surface and C/PEKK with prolonged dwell time at elevated temperature, a phenomenon that could not be explained by our current understanding. This work presents additional friction and differential scanning calorimetry experiments to demonstrate that the increased adhesion results from polymer-metal interactions that lead to polymer degradation and metal-polymer bonding. These findings underline the importance of tool material selection in both processing and characterization of tool-ply friction.

**Introduction**

High friction or adhesion between tool surface and composite material can lead to fiber waviness defects during the manufacture of thermoplastic composite parts [1–3]. Many research groups have characterized tool-ply friction under processing conditions in an effort to understand and prevent such defect formation. The characterization experiments involve the controlled sliding of a molten thermoplastic composite ply against a representative tool surface [4]. The measured shear stress at the interface typically displays a peak, followed by a steady-state plateau, where the peak and steady-state stress levels depend on temperature, applied normal pressure and sliding rate [5–7].

The mechanisms underlying the observed shear stress response are schematically illustrated in Fig. 1, and include friction, physical and/or chemical bonding, and viscoelastic deformation of the molten polymer. The peak shear stress could be caused by an adhesive bond or static friction between the fibers and the tool, which needs to be overcome before sliding can commence [8, 9]. Additionally, a thin polymer interlayer exists between the fibers and the metal surface, acting as a lubricating film. Here, the evolution of the shear stress is governed by the deformation of the matrix interlayer. The peak stress can then be caused by the nonlinear viscoelasticity of the matrix material [6, 7], a change in the interlayer thickness distribution [5], or by the development of slip at the tool-polymer [5] and/or the fiber-polymer interface [10, 11].

Although the current understanding of these phenomena is well-developed, our previous research [12] brought to light a curious phenomenon: the peak friction between mild steel and Solvay unidirectional C/PEKK significantly increased with dwell time at elevated temperature; an effect that was not observed for ply-ply friction. This unexpected finding challenged our understanding of the governing mechanisms and prompted additional experiments to identify the underlying causes. In this paper, we report results indicating that the increased peak friction is caused by strong polymer-steel bonding and polymer cross-linking, with the latter accelerated by

<figure><img src="image_486.png" /></figure>
<figure><img src="image_476.png" /><figcaption>Figure 1. Schematic illustration of the tool-ply interphase, highlighting the mechanisms that influence the friction response.</figcaption></figure>

contact with the mild steel. These findings have implications for tool material choices in manufacturing and friction characterization.

This paper is structured as follows. First, the materials and methods used are described, followed by a brief presentation of the results. In the discussion, we will interpret our findings within the context of the described mechanisms, while the final section presents the conclusions.

## Materials and methods

*Materials.* All experiments were conducted on APC C/PEKK unidirectional tape material from Solvay. This tape has a fiber volume fraction of 59% and a thickness of approximately 0.15 mm, while it exhibits a glass transition temperature of 159 °C and a melting temperature of 337 °C. A few experiments featured a 50 µm thick PEKK film, which according to the supplier has properties comparable to the matrix used in the APC tape.

The tool surface was represented by a non-alloy cold-rolled steel, known as DC01, with a chemical composition as provided in Table 1. The steel was supplied by Georg Martin GmbH as a foil with a thickness of 50 µm. Prior to the experimental work, the foil surface underwent cleaning using isopropanol. In certain experiments, we applied a release agent (Marbocote 227CEE) to the metal foils following the supplier's recommendation, leaving it to cure for 24 hours at room temperature. We also examined the friction between the composite tape material and polyimide films, which are commonly used in industry for facilitating part release. The polyimide film in question has a thickness of 25 µm, is known as Upilex 25S, and was supplied by UBE GmbH.

<table><caption>Table 1. Chemical composition (max weight %) of DC01 steel.</caption><thead><tr><th>C</th><th>Mn</th><th>P</th><th>S</th></tr></thead><tbody><tr><td>0.12</td><td>0.6</td><td>0.045</td><td>0.045</td></tr></tbody></table>

*Tool-Ply Friction Measurements.* The friction between the molten UD tape material and the metal foil or polyimide film was characterized using a well-established friction tester [4]. A tool-ply friction specimen comprises a UD tape, having a width of 50 mm, sandwiched between two metal foils or polyimide films with a width of 55 mm. The overlap between tape and metal foil or polyimide film was 65 mm. We considered the following contact pairs:

- APC vs. mild steel with and without release coating

- APC vs. polyimide film

- APC vs. mild steel with additional PEKK film at the interface

Fig. 2 schematically illustrates the working principle of the friction tester [13]. The tool-ply specimen is placed between two heated press platens, each measuring 50 × 50 mm², which apply a normal pressure. The ends of the metal foils are clamped, while the UD tape is pulled through and forced to slide against the metal foils. During the measurement, we logged the displacement and the required pull force. The sliding action was terminated after a displacement of approximately 10 mm. The pull force *F* was used to calculate the shear stress at the tool-ply interface as:
<figure><img src="image_487.png" /><figcaption>Figure 1. Schematic illustration of the tool-ply interphase, highlighting the mechanisms that influence the friction response.</figcaption></figure>

contact with the mild steel. These findings have implications for tool material choices in manufacturing and friction characterization.

This paper is structured as follows. First, the materials and methods used are described, followed by a brief presentation of the results. In the discussion, we will interpret our findings within the context of the described mechanisms, while the final section presents the conclusions.

## Materials and methods

*Materials.* All experiments were conducted on APC C/PEKK unidirectional tape material from Solvay. This tape has a fiber volume fraction of 59% and a thickness of approximately 0.15 mm, while it exhibits a glass transition temperature of 159 °C and a melting temperature of 337 °C. A few experiments featured a 50 µm thick PEKK film, which according to the supplier has properties comparable to the matrix used in the APC tape.

The tool surface was represented by a non-alloy cold-rolled steel, known as DC01, with a chemical composition as provided in Table 1. The steel was supplied by Georg Martin GmbH as a foil with a thickness of 50 µm. Prior to the experimental work, the foil surface underwent cleaning using isopropanol. In certain experiments, we applied a release agent (Marbocote 227CEE) to the metal foils following the supplier's recommendation, leaving it to cure for 24 hours at room temperature. We also examined the friction between the composite tape material and polyimide films, which are commonly used in industry for facilitating part release. The polyimide film in question has a thickness of 25 µm, is known as Upilex 25S, and was supplied by UBE GmbH.

<table><caption>Table 1. Chemical composition (max weight %) of DC01 steel.</caption><thead><tr><th>C</th><th>Mn</th><th>P</th><th>S</th></tr></thead><tbody><tr><td>0.12</td><td>0.6</td><td>0.045</td><td>0.045</td></tr></tbody></table>

*Tool-Ply Friction Measurements.* The friction between the molten UD tape material and the metal foil or polyimide film was characterized using a well-established friction tester [4]. A tool-ply friction specimen comprises a UD tape, having a width of 50 mm, sandwiched between two metal foils or polyimide films with a width of 55 mm. The overlap between tape and metal foil or polyimide film was 65 mm. We considered the following contact pairs:

- APC vs. mild steel with and without release coating

- APC vs. polyimide film

- APC vs. mild steel with additional PEKK film at the interface

Fig. 2 schematically illustrates the working principle of the friction tester [13]. The tool-ply specimen is placed between two heated press platens, each measuring 50 × 50 mm², which apply a normal pressure. The ends of the metal foils are clamped, while the UD tape is pulled through and forced to slide against the metal foils. During the measurement, we logged the displacement and the required pull force. The sliding action was terminated after a displacement of approximately 10 mm. The pull force *F* was used to calculate the shear stress at the tool-ply interface as:
<figure><img src="image_477.png" /><figcaption>Figure 2. Schematic illustration of the friction test.</figcaption></figure>

$$ \tau = \frac{F}{2A'} \qquad (1) $$

with A as the contact area.

In all friction experiments, we applied a normal pressure of 15 kPa, a sliding rate of 25 mm/min and a platen temperature of 385 °C. The dwell time between closing of the pressure platens and initiating sliding motion was varied between 2.5, 5, and 10 minutes.

*Differential Scanning Calorimetry.* We conducted differential scanning calorimetry (DSC) experiments, using a TA Instruments Discovery 250, to assess potential chemical changes in the tape materials resulting from the friction experiments. To this end, small pieces of tape material, with a total mass of 8 to 10 mg, were extracted from the APC friction specimens after testing. These DSC samples were then subjected to the following thermal cycle in a nitrogen atmosphere:

1. The samples were heated to 385 °C a rate of 20 °C/min;

2. Following the heating phase, an isothermal segment was maintained for 5 minutes;

3. Finally, the samples were cooled back to room temperature at a rate of 20 °C/min.

Using a heat of fusion of 130 J/g for the matrix, we determined the degree of crystallinity from the crystallization peak measured in the cooling traces. The degradation of PAEKs involves chain scission at the ether or carbonyl linkage, followed by branching or cross-linking [14, 15]. These latter two processes reduce chain mobility, increasing the polymer's viscosity and inhibiting its ability to crystallize [15]. In this work, we utilized the degree of crystallinity as an indicator for polymer degradation.

Differential scanning calorimetry experiments were also conducted on the as-received PEKK film and APC tape material. Four sample pans were prepared, each containing a different combination: one with small pieces of PEKK film, another with a mix of PEKK film and metal foil, a third with pieces of the APC tape, and a fourth with a mix of the APC tape and metal foil pieces. The goal of these experiments was to investigate the influence of the metal foil on the degradation of the matrix material. The samples were first subjected to a 10-minute dwell at elevated temperature under nitrogen atmosphere in the DSC equipment, followed by the same analysis step as outlined above.

*Scanning Electron Microscopy.* Scanning electron microscopy (SEM) images were captured to examine the bonding between the APC tape and the metal foil after a 10-minute dwell. The metal foil was peeled by hand from the APC tape, requiring substantial force. Subsequently, SEM images were obtained using a Jeol 7200F scanning electron microscope. Additionally, the tape surface was analyzed using Energy Dispersive X-ray (EDX) spectroscopy, using detectors from Oxford Instruments, operating at an acceleration voltage of 10 kV.

## Results

*Tool-Ply Friction.* Fig. 3 illustrates the interfacial shear stress evolution as a function of displacement for all contact pairs considered. The curves for APC in contact with the mild steel foils, shown in Fig. 3a-b, display an initial peak at the start of the test, followed by a subsequent
<figure><img src="image_488.png" /><figcaption>Figure 2. Schematic illustration of the friction test.</figcaption></figure>

$$ \tau = \frac{F}{2A'} \qquad (1) $$

with A as the contact area.

In all friction experiments, we applied a normal pressure of 15 kPa, a sliding rate of 25 mm/min and a platen temperature of 385 °C. The dwell time between closing of the pressure platens and initiating sliding motion was varied between 2.5, 5, and 10 minutes.

*Differential Scanning Calorimetry.* We conducted differential scanning calorimetry (DSC) experiments, using a TA Instruments Discovery 250, to assess potential chemical changes in the tape materials resulting from the friction experiments. To this end, small pieces of tape material, with a total mass of 8 to 10 mg, were extracted from the APC friction specimens after testing. These DSC samples were then subjected to the following thermal cycle in a nitrogen atmosphere:

1. The samples were heated to 385 °C a rate of 20 °C/min;

2. Following the heating phase, an isothermal segment was maintained for 5 minutes;

3. Finally, the samples were cooled back to room temperature at a rate of 20 °C/min.

Using a heat of fusion of 130 J/g for the matrix, we determined the degree of crystallinity from the crystallization peak measured in the cooling traces. The degradation of PAEKs involves chain scission at the ether or carbonyl linkage, followed by branching or cross-linking [14, 15]. These latter two processes reduce chain mobility, increasing the polymer's viscosity and inhibiting its ability to crystallize [15]. In this work, we utilized the degree of crystallinity as an indicator for polymer degradation.

Differential scanning calorimetry experiments were also conducted on the as-received PEKK film and APC tape material. Four sample pans were prepared, each containing a different combination: one with small pieces of PEKK film, another with a mix of PEKK film and metal foil, a third with pieces of the APC tape, and a fourth with a mix of the APC tape and metal foil pieces. The goal of these experiments was to investigate the influence of the metal foil on the degradation of the matrix material. The samples were first subjected to a 10-minute dwell at elevated temperature under nitrogen atmosphere in the DSC equipment, followed by the same analysis step as outlined above.

*Scanning Electron Microscopy.* Scanning electron microscopy (SEM) images were captured to examine the bonding between the APC tape and the metal foil after a 10-minute dwell. The metal foil was peeled by hand from the APC tape, requiring substantial force. Subsequently, SEM images were obtained using a Jeol 7200F scanning electron microscope. Additionally, the tape surface was analyzed using Energy Dispersive X-ray (EDX) spectroscopy, using detectors from Oxford Instruments, operating at an acceleration voltage of 10 kV.

## Results

*Tool-Ply Friction.* Fig. 3 illustrates the interfacial shear stress evolution as a function of displacement for all contact pairs considered. The curves for APC in contact with the mild steel foils, shown in Fig. 3a-b, display an initial peak at the start of the test, followed by a subsequent
decrease to a plateau value. Notably, the peak shear stress increases with longer dwell times, reaching levels exceeding the maximum measurable capacity of 200 kPa for metal foils without release agent. Importantly, these results align with our earlier findings [12]. In contrast, the friction response for APC against the polyimide release film, shown in Fig. 3c, does not exhibit an initial peak, and the overall stress evolution also does not show any significant dependency on dwell time.

Fig. 3d shows the shear stress evolution for APC against steel with the PEKK film inserted at the tool-ply interface. Compared to the corresponding case without film, shown in Fig. 3a, both the peak and steady-state stress levels have reduced significantly. Further, although the peak stress increases with dwell time, the increase is not as dramatic as observed in the case without PEKK film, while the steady-state shear stress now also shows an increase with dwell time.

*Differential Scanning Calorimetry.* The extended dwell at high temperature may have induced chemical changes in the polymer. To investigate this hypothesis, tape samples were extracted from the APC friction specimens and subjected to a DSC analysis in a nitrogen atmosphere. The cooling trace was analyzed to obtain the degree of crystallinity. The results are plotted in the left graph in Fig. 4, which shows the calculated degree of crystallinity as a function of dwell time for different tool contacts. These results indicate a reduced ability to crystallize with increasing dwell time for the APC tape specimens in contact with the metal foil. In contrast, this reduced ability to crystallize was not observed for specimens extracted from friction specimens concerning the release-coated steel or the polyimide film.

<figure><img src="image_478.png" /><figcaption>Figure 3. Interfacial shear stress evolution for APC against various tool surfaces, where '+ RC' indicates that a release coating was applied to the steel surface. The colors in the graphs represent different dwell times between closing of the pressure platens and initiating sliding action. All experiments were performed at a temperature of 385 °C, with a normal pressure of 15 kPa and a sliding rate of 25 mm/min.</figcaption></figure>
decrease to a plateau value. Notably, the peak shear stress increases with longer dwell times, reaching levels exceeding the maximum measurable capacity of 200 kPa for metal foils without release agent. Importantly, these results align with our earlier findings [12]. In contrast, the friction response for APC against the polyimide release film, shown in Fig. 3c, does not exhibit an initial peak, and the overall stress evolution also does not show any significant dependency on dwell time.

Fig. 3d shows the shear stress evolution for APC against steel with the PEKK film inserted at the tool-ply interface. Compared to the corresponding case without film, shown in Fig. 3a, both the peak and steady-state stress levels have reduced significantly. Further, although the peak stress increases with dwell time, the increase is not as dramatic as observed in the case without PEKK film, while the steady-state shear stress now also shows an increase with dwell time.

*Differential Scanning Calorimetry.* The extended dwell at high temperature may have induced chemical changes in the polymer. To investigate this hypothesis, tape samples were extracted from the APC friction specimens and subjected to a DSC analysis in a nitrogen atmosphere. The cooling trace was analyzed to obtain the degree of crystallinity. The results are plotted in the left graph in Fig. 4, which shows the calculated degree of crystallinity as a function of dwell time for different tool contacts. These results indicate a reduced ability to crystallize with increasing dwell time for the APC tape specimens in contact with the metal foil. In contrast, this reduced ability to crystallize was not observed for specimens extracted from friction specimens concerning the release-coated steel or the polyimide film.

<figure><img src="image_489.png" /><figcaption>Figure 3. Interfacial shear stress evolution for APC against various tool surfaces, where '+ RC' indicates that a release coating was applied to the steel surface. The colors in the graphs represent different dwell times between closing of the pressure platens and initiating sliding action. All experiments were performed at a temperature of 385 °C, with a normal pressure of 15 kPa and a sliding rate of 25 mm/min.</figcaption></figure>
<figure><img src="image_479.png" /><figcaption>Figure 4. Left: Degree of crystallinity, after cooling at 20 °C/min in a DSC, for APC tape samples extracted from the friction specimens as a function of the dwell time. The color of the markers represent the different tool surfaces that were in contact with the tape. The shaded gray area indicates the maximum and minimum degree crystallinity measured on the as-received, and untreated, APC tape without a dwell. Right: Degree of crystallinity, after cooling at 20 °C/min in a DSC, for as-received APC tape and PEKK film samples with and without metal.</figcaption></figure>

For reference, we also analyzed the as-received APC tape and PEKK film using DSC, both with and without small pieces of steel foil in the same pan. Prior to the analysis, the specimens were subjected to 10-minute dwell. The right graph in Fig. 4 shows the measured degrees of crystallinity. Overall, the results correspond with the data measured on the friction specimens, showing a decrease in the polymer's ability to crystallize when the APC tape or the PEKK film has been in contact with the metal film during a dwell in melt.

*Microscopy.* Fig. 5 presents two SEM images of a tape surface extracted from an APC-steel friction specimen subjected to a 10-minute dwell. As mentioned, significant force was required to peel the tape in question from the steel foil at room temperature. The images reveal the presence of small iron particles on the tape surface. EDX analysis confirmed that these particles contain a high fraction of iron, indicating material transfer from the steel foil surface to the ply.

SEM images of the corresponding steel foil surface are shown in Fig. 6. The horizontal lines on the metal surface are also present on the as-received foils and are oriented in the foil's rolling direction, which was perpendicular to the sliding direction used in the friction experiment. The SEM images illustrate damage in the form of craters on the foil's surface, which is in agreement with the iron particles found on the tape surface. These observations confirm a strong bond between the mild steel surface and the tape, after a 10-minute dwell at high temperature.

## Discussion

In this section, we will discuss the possible interactions between mild steel and the PEKK polymer in the APC tape. Focusing on metal-catalyzed branching or cross-linking and surface bonding, we propose a mechanistic model of the interphase region to explain our experimental findings.
<figure><img src="image_490.png" /><figcaption>**Figure 4.** Left: Degree of crystallinity, after cooling at 20 °C/min in a DSC, for APC tape samples extracted from the friction specimens as a function of the dwell time. The color of the markers represent the different tool surfaces that were in contact with the tape. The shaded gray area indicates the maximum and minimum degree crystallinity measured on the as-received, and untreated, APC tape without a dwell. Right: Degree of crystallinity, after cooling at 20 °C/min in a DSC, for as-received APC tape and PEKK film samples with and without metal.</figcaption></figure>

For reference, we also analyzed the as-received APC tape and PEKK film using DSC, both with and without small pieces of steel foil in the same pan. Prior to the analysis, the specimens were subjected to 10-minute dwell. The right graph in Fig. 4 shows the measured degrees of crystallinity. Overall, the results correspond with the data measured on the friction specimens, showing a decrease in the polymer's ability to crystallize when the APC tape or the PEKK film has been in contact with the metal film during a dwell in melt.

*Microscopy.* Fig. 5 presents two SEM images of a tape surface extracted from an APC-steel friction specimen subjected to a 10-minute dwell. As mentioned, significant force was required to peel the tape in question from the steel foil at room temperature. The images reveal the presence of small iron particles on the tape surface. EDX analysis confirmed that these particles contain a high fraction of iron, indicating material transfer from the steel foil surface to the ply.

SEM images of the corresponding steel foil surface are shown in Fig. 6. The horizontal lines on the metal surface are also present on the as-received foils and are oriented in the foil's rolling direction, which was perpendicular to the sliding direction used in the friction experiment. The SEM images illustrate damage in the form of craters on the foil's surface, which is in agreement with the iron particles found on the tape surface. These observations confirm a strong bond between the mild steel surface and the tape, after a 10-minute dwell at high temperature.

**Discussion**

In this section, we will discuss the possible interactions between mild steel and the PEKK polymer in the APC tape. Focusing on metal-catalyzed branching or cross-linking and surface bonding, we propose a mechanistic model of the interphase region to explain our experimental findings.
<figure><img src="image_480.png" /><figcaption>Figure 5. Scanning electron microscopy images showing the presence of iron particles on the fracture surface of the APC tape. The left image shows an overview, while the right a detailed view of an iron particle. Both inset images show the corresponding EDX maps for iron. The tape was extracted from an APC-metal friction specimen that was subjected to a 10-minute dwell.</figcaption></figure>

<figure><img src="image_481.png" /><figcaption>Figure 6. Scanning electron microscopy images showing damage on the metal surface. The metal foil was extracted from an APC-metal friction specimen that was subjected to a 10-minute dwell.</figcaption></figure>

**Metal-Catalyzed Chain Branching or Cross-Linking.** The DSC results reported in Fig. 4 show a reduction of the polymer's ability to crystallize after prolonged contact with mild steel at elevated temperature. This effect indicates chain branching or cross-linking [15], and is only observed in cases involving contact with the mild steel, suggesting that the steel accelerates or causes polymer degradation.

Literature on the catalytic effect of a metal tool surface on polymer degradation during processing is limited, and we did not find any information specific for poly(aryl-ether-ketones). Some studies on polycarbonates are available though, and provide valuable insight. Sonnenberg et al. [16] studied the interactions between tool steel and a polycarbonate melt. Through X-ray Photoelectron Spectroscopy and Auger Electron Spectroscopy, they demonstrated the diffusion of metal atoms from the steel surface into the polymer melt. Notably, iron ions reacted with the polymer chains, resulting in the formation of iron-polycarbonate complexes that agglomerate through cross-linking and lead to the immobilization of polymer chains. Further investigations by Theile-Rasche et al. [17] correlated the degree of degradation in polycarbonate with the percentage of iron at the alloy surface. Although degradation was most pronounced near the interface, their experiments showed a degraded layer with a thickness 250 to 300 µm after one hour dwell at processing temperature. Notably, these authors also found that the presence of chromium inhibits polymer degradation. Possibly this is caused by a chromium oxide layer that inhibits iron diffusion.

A similar branching or cross-linking mechanism may be active in PEKK when it is in contact with a mild steel surface. The presented friction data supports this idea. Most notably, the extreme peak shear stress observed in the case of APC tape and the mild steel indicates that the polymer
<figure><img src="image_491.png" /><figcaption>Figure 5. Scanning electron microscopy images showing the presence of iron particles on the fracture surface of the APC tape. The left image shows an overview, while the right a detailed view of an iron particle. Both inset images show the corresponding EDX maps for iron. The tape was extracted from an APC-metal friction specimen that was subjected to a 10-minute dwell.</figcaption></figure>

<figure><img src="image_492.png" /><figcaption>Figure 6. Scanning electron microscopy images showing damage on the metal surface. The metal foil was extracted from an APC-metal friction specimen that was subjected to a 10-minute dwell.</figcaption></figure>

**Metal-Catalyzed Chain Branching or Cross-Linking.** The DSC results reported in Fig. 4 show a reduction of the polymer's ability to crystallize after prolonged contact with mild steel at elevated temperature. This effect indicates chain branching or cross-linking [15], and is only observed in cases involving contact with the mild steel, suggesting that the steel accelerates or causes polymer degradation.

Literature on the catalytic effect of a metal tool surface on polymer degradation during processing is limited, and we did not find any information specific for poly(aryl-ether-ketones). Some studies on polycarbonates are available though, and provide valuable insight. Sonnenberg et al. [16] studied the interactions between tool steel and a polycarbonate melt. Through X-ray Photoelectron Spectroscopy and Auger Electron Spectroscopy, they demonstrated the diffusion of metal atoms from the steel surface into the polymer melt. Notably, iron ions reacted with the polymer chains, resulting in the formation of iron-polycarbonate complexes that agglomerate through cross-linking and lead to the immobilization of polymer chains. Further investigations by Theile-Rasche et al. [17] correlated the degree of degradation in polycarbonate with the percentage of iron at the alloy surface. Although degradation was most pronounced near the interface, their experiments showed a degraded layer with a thickness 250 to 300 µm after one hour dwell at processing temperature. Notably, these authors also found that the presence of chromium inhibits polymer degradation. Possibly this is caused by a chromium oxide layer that inhibits iron diffusion.

A similar branching or cross-linking mechanism may be active in PEKK when it is in contact with a mild steel surface. The presented friction data supports this idea. Most notably, the extreme peak shear stress observed in the case of APC tape and the mild steel indicates that the polymer
present between the fibers and the steel has cross-linked up to a degree that it acts as a gel or solid instead of a lubricating film. At shorter dwell times, branching and cross-linking may cause an increase in viscosity but not yet to an extent that it leads to complete immobilization of the polymer chains. Similarly, the increase in peak and steady-state shear stress with dwell time in the case of the inserted PEKK film between APC and the uncoated steel, illustrated in Fig. 3d, also point to an increased viscosity of the lubricating polymer film. Here, the film can still act as a lubricant, even after a 10-minute dwell, because of its substantial thickness.

Finally, although the DSC data did not indicate chain branching or cross-linking for the release-coated friction specimens, the increasing peak shear stress with dwell time, shown in 3b, suggests otherwise. Possibly, the release coating acts as a barrier, limiting the diffusion of iron into the polymer and minimizing branching or cross-linking to only a surface layer of the tape, where it can have a significant influence on the measured shear stress but has a negligible effect on the measured DSC response, as the latter is governed by the bulk of the tape.

*Surface Bonding.* The high shear stresses required to initiate sliding between the APC tape and the uncoated metal, particularly after a 5- or 10-minute dwell, suggests strong bonding between the polymer and the metal. This is supported by the fracture surfaces, presented in Fig. 5 and 6, which illustrate damage to the steel surface and transfer of metal particles to the tape. While the nature of this bonding is unclear, our findings point to a chemical interaction between the polymer and metal surfaces.

Such chemical interactions have been reported in literature, albeit for different material systems. In the late 70s and early 80s, Burkstrand [18–20] for example demonstrated increased polymer-metal adhesive strength that could be related to the formation of metal-oxygen-polymer complexes at the interface for several polymers including polystyrene and PMMA. Recently, Wiesing et al. [21] employed XPS analysis to show the formation of Fe(II)-rich oxides at the interface between tool steel and molten polycarbonate.

The observed high peak stresses also imply strong bonding between fiber and matrix. In ply- ply friction experiments, we found an apparent critical fiber-matrix slip stress of approximately 35 and 55 kPa for C/PEEK and C/LM-PAEK, respectively [10, 11]. The proposed slip mechanism in these cases involves the disentanglement of chains, that are adsorbed on the fiber, from the chains in the bulk. Our inability to induce any sliding action in the APC-steel case, with a dwell of 10 minutes, suggests that polymer disentanglement is inhibited, which is consistent with the branching and cross-linking deduced from the DSC experiments.

*Proposed Mechanism.* The observed increase in friction between mild steel and APC with prolonged dwell time can be attributed to polymer-metal interactions. Based on our experimental results and existing literature, we propose a mechanistic model, schematically illustrated in Fig. 7. First, iron ions diffuse into the molten PEKK polymer at the interphase between fibers and tool. These ions interact with the polymer chains, leading to metal-polymer complexes through branching or cross-linking, similar as was observed for polycarbonates [16, 17]. As a result, the polymer viscosity increases up to a point where the thin layer between the metal and fibers becomes a gel and loses its lubricating function. Simultaneously, chemical bonding takes place between the polymer and the steel, possibly by the formation of metal-oxygen-polymer complexes at the interface [18–21]. The resulting adhesion between polymer is so strong that excessive stress is required to initiate sliding between tape and metal.
present between the fibers and the steel has cross-linked up to a degree that it acts as a gel or solid instead of a lubricating film. At shorter dwell times, branching and cross-linking may cause an increase in viscosity but not yet to an extent that it leads to complete immobilization of the polymer chains. Similarly, the increase in peak and steady-state shear stress with dwell time in the case of the inserted PEKK film between APC and the uncoated steel, illustrated in Fig. 3d, also point to an increased viscosity of the lubricating polymer film. Here, the film can still act as a lubricant, even after a 10-minute dwell, because of its substantial thickness.

Finally, although the DSC data did not indicate chain branching or cross-linking for the release-coated friction specimens, the increasing peak shear stress with dwell time, shown in 3b, suggests otherwise. Possibly, the release coating acts as a barrier, limiting the diffusion of iron into the polymer and minimizing branching or cross-linking to only a surface layer of the tape, where it can have a significant influence on the measured shear stress but has a negligible effect on the measured DSC response, as the latter is governed by the bulk of the tape.

*Surface Bonding.* The high shear stresses required to initiate sliding between the APC tape and the uncoated metal, particularly after a 5- or 10-minute dwell, suggests strong bonding between the polymer and the metal. This is supported by the fracture surfaces, presented in Fig. 5 and 6, which illustrate damage to the steel surface and transfer of metal particles to the tape. While the nature of this bonding is unclear, our findings point to a chemical interaction between the polymer and metal surfaces.

Such chemical interactions have been reported in literature, albeit for different material systems. In the late 70s and early 80s, Burkstrand [18–20] for example demonstrated increased polymer-metal adhesive strength that could be related to the formation of metal-oxygen-polymer complexes at the interface for several polymers including polystyrene and PMMA. Recently, Wiesing et al. [21] employed XPS analysis to show the formation of Fe(II)-rich oxides at the interface between tool steel and molten polycarbonate.

The observed high peak stresses also imply strong bonding between fiber and matrix. In ply- ply friction experiments, we found an apparent critical fiber-matrix slip stress of approximately 35 and 55 kPa for C/PEEK and C/LM-PAEK, respectively [10, 11]. The proposed slip mechanism in these cases involves the disentanglement of chains, that are adsorbed on the fiber, from the chains in the bulk. Our inability to induce any sliding action in the APC-steel case, with a dwell of 10 minutes, suggests that polymer disentanglement is inhibited, which is consistent with the branching and cross-linking deduced from the DSC experiments.

*Proposed Mechanism.* The observed increase in friction between mild steel and APC with prolonged dwell time can be attributed to polymer-metal interactions. Based on our experimental results and existing literature, we propose a mechanistic model, schematically illustrated in Fig. 7. First, iron ions diffuse into the molten PEKK polymer at the interphase between fibers and tool. These ions interact with the polymer chains, leading to metal-polymer complexes through branching or cross-linking, similar as was observed for polycarbonates [16, 17]. As a result, the polymer viscosity increases up to a point where the thin layer between the metal and fibers becomes a gel and loses its lubricating function. Simultaneously, chemical bonding takes place between the polymer and the steel, possibly by the formation of metal-oxygen-polymer complexes at the interface [18–21]. The resulting adhesion between polymer is so strong that excessive stress is required to initiate sliding between tape and metal.
<figure><img src="image_482.png" /><figcaption>Figure 7. Schematic illustration of the tool-ply interphase, highlighting the mechanisms that govern the observed peak shear stress for APC and steel (inspired by [22]).</figcaption></figure>

Theile-Rasche et al. demonstrated that an increase in chromium content in the steel can strongly inhibit polymer degradation [17]. To validate this finding, we conducted additional friction experiments using stainless steel foils, from grade X5CrNi18-10, featuring an 18% chromium content (see Table 2). Fig. 8 shows the shear stress evolution for the different dwell times, with the response for mild steel included for direct comparison. Some initial slip was observed between the stainless steel foil and the pressure platens, noticeable at the beginning of the curve (< 1 mm). In contrast to mild steel, where the peak shear stress increases quickly with dwell time, the influence of dwell time on shear stress evolution for stainless steel is minimal. This observation aligns with the work of Theile-Rasche et al. [17] and supports the hypothesis that iron ions play a key role in polymer degradation. Overall, these results underline the importance of tool material selection on friction and, hence, defect formation during manufacturing.

<table><caption>Table 2. Chemical composition (max weight %) of X5CrNi18-10 stainless steel.</caption><thead><tr><th>C</th><th>Si</th><th>Mn</th><th>Ni</th><th>P</th><th>S</th><th>Cr</th><th>N</th></tr></thead><tbody><tr><td>0.07</td><td>1</td><td>2</td><td>8-10.5</td><td>0.045</td><td>0.015</td><td>17.5-19.5</td><td>0.11</td></tr></tbody></table>

<figure><img src="image_483.png" /><figcaption>Figure 8. Interfacial shear stress evolution for APC against mild steel (DC01) and stainless steel (X5CrNi18-10). The colors in the graphs represent different dwell times between closing of the pressure platens and initiating sliding action. All experiments were performed at a temperature of 385 °C, with a normal pressure of 15 kPa and a sliding rate of 25 mm/min.</figcaption></figure>

## Conclusions

Investigating the impact of dwell time on tool-ply friction is essential, as excessive friction can cause defect formation during the manufacturing of thermoplastic composite parts. In this work,
<figure><img src="image_493.png" /><figcaption>Figure 7. Schematic illustration of the tool-ply interphase, highlighting the mechanisms that govern the observed peak shear stress for APC and steel (inspired by [22]).</figcaption></figure>

Theile-Rasche et al. demonstrated that an increase in chromium content in the steel can strongly inhibit polymer degradation [17]. To validate this finding, we conducted additional friction experiments using stainless steel foils, from grade X5CrNi18-10, featuring an 18% chromium content (see Table 2). Fig. 8 shows the shear stress evolution for the different dwell times, with the response for mild steel included for direct comparison. Some initial slip was observed between the stainless steel foil and the pressure platens, noticeable at the beginning of the curve (< 1 mm). In contrast to mild steel, where the peak shear stress increases quickly with dwell time, the influence of dwell time on shear stress evolution for stainless steel is minimal. This observation aligns with the work of Theile-Rasche et al. [17] and supports the hypothesis that iron ions play a key role in polymer degradation. Overall, these results underline the importance of tool material selection on friction and, hence, defect formation during manufacturing.

<table><caption>Table 2. Chemical composition (max weight %) of X5CrNi18-10 stainless steel.</caption><thead><tr><th>C</th><th>Si</th><th>Mn</th><th>Ni</th><th>P</th><th>S</th><th>Cr</th><th>N</th></tr></thead><tbody><tr><td>0.07</td><td>1</td><td>2</td><td>8-10.5</td><td>0.045</td><td>0.015</td><td>17.5-19.5</td><td>0.11</td></tr></tbody></table>

<figure><img src="image_494.png" /><figcaption>Figure 8. Interfacial shear stress evolution for APC against mild steel (DC01) and stainless steel (X5CrNi18-10). The colors in the graphs represent different dwell times between closing of the pressure platens and initiating sliding action. All experiments were performed at a temperature of 385 °C, with a normal pressure of 15 kPa and a sliding rate of 25 mm/min.</figcaption></figure>

## Conclusions

Investigating the impact of dwell time on tool-ply friction is essential, as excessive friction can cause defect formation during the manufacturing of thermoplastic composite parts. In this work,
we investigated the effect of dwell time on the tool-ply friction response in melt for APC tape materials. The tool materials included mild steel, stainless steel and polyimide film. In line with earlier observations, the peak shear stress was found to increase significantly with dwell time in the case of the uncoated mild steel surface; an effect which was not observed for stainless steel or polyimide film.

The observed increase in peak shear stress for the mild steel surface is believed to be caused by metal-polymer interactions. Iron diffuses from the surface into the polymer, where it forms iron-polymer complexes that cause polymer chain branching and cross-linking. As a consequence the viscosity of the polymer increases up to a point where it cannot act anymore as a lubricating film. Simultaneously, these interactions also lead to strong adhesion between the polymer and the steel surface. Combined, these two effects result in strong adhesion between the tape and mild steel. The minimal impact of dwell time on the shear stress evolution in the case of a stainless steel surface, in contrast to mild steel, aligns with literature and emphasizes the role of iron ions in polymer chain branching and cross-linking. Overall, our findings emphasize the importance of tool material selection in manufacturing processes to ensure defect-free thermoplastic composite parts.

## Acknowledgements

This work was performed as part of the MaterialenNL research program under project number 17880, which is (partly) financed by the Dutch Research Council (NWO). The authors also gratefully acknowledge the financial and technical support from the industrial and academic members of the ThermoPlastic composites Research Center (TPRC).

## Data availability

The data belonging to this work is available via the 4TU.ResearchData repository [23].

## References

[1] D. Kugler and T.J. Moon, Identification of the most significant processing parameters on the development of fiber waviness in thin laminates, J. Compos. Mater. 36:12 (2002) 1451-1479.
https://doi.org/10.1177/0021998302036012575

[2] M. Thor, M.G.R. Sause, R.M. Hinterhölzl, Mechanisms of origin and classification of out-of-plane fiber waviness in composite materials--A review, J. Compos. Sci. 4:3 (2020) p.130.
https://doi.org/10.3390/jcs4030130

[3] E.T.M Krämer, W.J.B. Grouve, L.L. Warnet, S. Koussios, R. Akkerman, Tool-ply interaction in the formation of waviness during C/PEEK consolidation, Compos. - A 144 (2021) p. 106327.
https://doi.org/10.1016/j.compositesa.2021.106327

[4] U. Sachs et al., Characterization of the dynamic friction of woven fabrics: Experimental methods and benchmark results, Compos. - A: Appl. Sci. Manuf. 67 (2014) 289-298.
https://doi.org/10.1016/j.compositesa.2014.08.026

[5] A.M. Murtagh, J.J. Lennon, and P.J. Mallon, Surface friction effects related to pressforming of continuous fibre thermoplastic composites, Compos. Manuf. 6:34 (1995) 169-175.
https://doi.org/10.1016/0956-7143(95)95008-M

[6] R.H.W. Ten Thije, R. Akkerman, M. Ubbink, and L. van der Meer, A lubrication approach to friction in thermoplastic composites forming processes, Compos. – A 42:8 (2011) 950-960.
https://doi.org/10.1016/j.compositesa.2011.03.023

[7] A. Kapshammer, D. Laresser, M.C. Miron, F. Baudach, and Z. Major, Characterization and modeling of ply/tool and ply/ply slippage phenomena of unidirectional polycarbonate CF tapes, Polymers 15:17 (2023) p. 3520.
https://doi.org/10.3390/polym15173520

[8] S.R. Morris and C. Sun, An investigation of interply slip behaviour in AS4/PEEK at forming temperatures, Compos. Manuf. 5:4 (1994) 217-224.
https://doi.org/10.1016/0956-7143(94)90136-8
we investigated the effect of dwell time on the tool-ply friction response in melt for APC tape materials. The tool materials included mild steel, stainless steel and polyimide film. In line with earlier observations, the peak shear stress was found to increase significantly with dwell time in the case of the uncoated mild steel surface; an effect which was not observed for stainless steel or polyimide film.

The observed increase in peak shear stress for the mild steel surface is believed to be caused by metal-polymer interactions. Iron diffuses from the surface into the polymer, where it forms iron-polymer complexes that cause polymer chain branching and cross-linking. As a consequence the viscosity of the polymer increases up to a point where it cannot act anymore as a lubricating film. Simultaneously, these interactions also lead to strong adhesion between the polymer and the steel surface. Combined, these two effects result in strong adhesion between the tape and mild steel. The minimal impact of dwell time on the shear stress evolution in the case of a stainless steel surface, in contrast to mild steel, aligns with literature and emphasizes the role of iron ions in polymer chain branching and cross-linking. Overall, our findings emphasize the importance of tool material selection in manufacturing processes to ensure defect-free thermoplastic composite parts.

## Acknowledgements

This work was performed as part of the MaterialenNL research program under project number 17880, which is (partly) financed by the Dutch Research Council (NWO). The authors also gratefully acknowledge the financial and technical support from the industrial and academic members of the ThermoPlastic composites Research Center (TPRC).

## Data availability

The data belonging to this work is available via the 4TU.ResearchData repository [23].

## References

[1] D. Kugler and T.J. Moon, Identification of the most significant processing parameters on the development of fiber waviness in thin laminates, J. Compos. Mater. 36:12 (2002) 1451-1479.
https://doi.org/10.1177/0021998302036012575

[2] M. Thor, M.G.R. Sause, R.M. Hinterhölzl, Mechanisms of origin and classification of out-of-plane fiber waviness in composite materials--A review, J. Compos. Sci. 4:3 (2020) p.130.
https://doi.org/10.3390/jcs4030130

[3] E.T.M Krämer, W.J.B. Grouve, L.L. Warnet, S. Koussios, R. Akkerman, Tool-ply interaction in the formation of waviness during C/PEEK consolidation, Compos. - A 144 (2021) p. 106327.
https://doi.org/10.1016/j.compositesa.2021.106327

[4] U. Sachs et al., Characterization of the dynamic friction of woven fabrics: Experimental methods and benchmark results, Compos. - A: Appl. Sci. Manuf. 67 (2014) 289-298.
https://doi.org/10.1016/j.compositesa.2014.08.026

[5] A.M. Murtagh, J.J. Lennon, and P.J. Mallon, Surface friction effects related to pressforming of continuous fibre thermoplastic composites, Compos. Manuf. 6:34 (1995) 169-175.
https://doi.org/10.1016/0956-7143(95)95008-M

[6] R.H.W. Ten Thije, R. Akkerman, M. Ubbink, and L. van der Meer, A lubrication approach to friction in thermoplastic composites forming processes, Compos. – A 42:8 (2011) 950-960.
https://doi.org/10.1016/j.compositesa.2011.03.023

[7] A. Kapshammer, D. Laresser, M.C. Miron, F. Baudach, and Z. Major, Characterization and modeling of ply/tool and ply/ply slippage phenomena of unidirectional polycarbonate CF tapes, Polymers 15:17 (2023) p. 3520.
https://doi.org/10.3390/polym15173520

[8] S.R. Morris and C. Sun, An investigation of interply slip behaviour in AS4/PEEK at forming temperatures, Compos. Manuf. 5:4 (1994) 217-224.
https://doi.org/10.1016/0956-7143(94)90136-8
[9] Y. Zhao, T. Zhang, H. Li, and B. Zhang, Characterization of prepreg-prepreg and prepreg-tool friction for unidirectional carbon fiber/epoxy prepreg during hot diaphragm forming process, Polym. Test 84 (2020) p. 106440. https://doi.org/10.1016/j.polymertesting.2020.106440

[10] E.R. Pierik, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, Prediction of the peak and steady-state ply-ply friction response for UD C/PAEK tapes, Compos. - A 163 (2022) p. 107185. https://doi.org/10.1016/j.compositesa.2022.107185

[11] E.R. Pierik, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, Modeling the effect of temperature and pressure on the peak and steady-state ply-ply friction response for UD C/PAEK tapes," Compos. - A 173 (2023) p. 107671. https://doi.org/10.1016/j.compositesa.2023.107671

[12] E.R. Pierik, J. Liddiard, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, On the effect of release agent and heating time on tool-ply friction of thermoplastic composite in melt, Key Engineering Materials 926 (2022) 1276-1283. https://doi.org/10.4028/p-e6m2a7

[13] R.H.W. ten Thije and R. Akkerman, Design of an experimental setup to measure tool-ply and ply-ply friction in thermoplastic laminates, Int. J. Mater. Form. 2:1 (2009) 197-200. https://doi.org/10.1007/s12289-009-0638-y

[14] P. Patel, T.R. Hull, R.W. McCabe, D. Flath, J. Grasmeder, and M. Percy, Mechanism of thermal decomposition of poly(ether ether ketone) (PEEK) from a review of decomposition studies, Polym. Degrad. Stab. 95:5 (2010) 709-718. https://doi.org/10.1016/j.polymdegradstab.2010.01.024

[15] C. Croshaw, L. Hamernik, L. Ghanbari, A. Browning, and J. Wiggins, Melt-state degradation mechanism of poly (ether ketone ketone): the role of branching on crystallization and rheological behavior, Polym. Degrad. Stab. 200 (2022) p. 109968. https://doi.org/10.1016/j.polymdegradstab.2022.109968

[16] M. Sonnenberg et al., Polymer-induced metal diffusion during plastic processing: a reason for deposit formation, J. Polym. Eng. 39:5 (2019) 472-480. https://doi.org/10.1515/polyeng-2018-0324

[17] C. Theile-Rasche, T. Meng, T. de los Arcos, and G. Grundmeier, "Analysis of polycarbonate degradation at melt/FeCr-alloy interfaces as a function of the alloy composition by means of combinatorial thin film chemistry," SN Appl. Sci. 5:10 (2023). https://doi.org/10.1007/s42452-023-05441-5

[18] J.M. Burkstrand, Formation of a copper-oxygen-polymer complex on polystyrene, Appl. Phys. Lett. 33:5 (1978) 387-389. https://doi.org/10.1063/1.90405

[19] J.M. Burkstrand, "Substrate effects on the electronic structure of metal overlayers--an XPS study of polymer-metal interfaces," Phys. Rev. B. 20:12 (1979) 4853-4858. https://doi.org/10.1103/PhysRevB.20.4853

[20] J.M. Burkstrand, "Metal-polymer interfaces: Adhesion and X-ray photoemission studies," J. Appl. Phys. 52:7 (1981) 4795-4800. https://doi.org/10.1063/1.329320

[21] M. Wiesing, T. de los Arcos, M. to Baben, H. Rueß, J. Schneider, and G. Grundmeier, Analysis of the inhibition of thermal degradation of molten polycarbonate at tool steel interfaces by thin TiAlN coatings, Polym. Degrad. Stab. 143 (2017) 196-206. https://doi.org/10.1016/j.polymdegradstab.2017.07.013

[22] G. Raos and B. Zappone, Polymer adhesion: Seeking new solutions for an old problem, Macromolecules 54:23 (2021) 10617-10644. https://doi.org/10.1021/acs.macromol.1c01182

[23] Data for this publication: https://doi.org/10.4121/f2a26463-a241-4698-bb3d-de31657d41d3
[9] Y. Zhao, T. Zhang, H. Li, and B. Zhang, Characterization of prepreg-prepreg and prepreg-tool friction for unidirectional carbon fiber/epoxy prepreg during hot diaphragm forming process, Polym. Test 84 (2020) p. 106440. https://doi.org/10.1016/j.polymertesting.2020.106440

[10] E.R. Pierik, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, Prediction of the peak and steady-state ply-ply friction response for UD C/PAEK tapes, Compos. - A 163 (2022) p. 107185. https://doi.org/10.1016/j.compositesa.2022.107185

[11] E.R. Pierik, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, Modeling the effect of temperature and pressure on the peak and steady-state ply-ply friction response for UD C/PAEK tapes," Compos. - A 173 (2023) p. 107671. https://doi.org/10.1016/j.compositesa.2023.107671

[12] E.R. Pierik, J. Liddiard, W.J.B. Grouve, S. Wijskamp, and R. Akkerman, On the effect of release agent and heating time on tool-ply friction of thermoplastic composite in melt, Key Engineering Materials 926 (2022) 1276-1283. https://doi.org/10.4028/p-e6m2a7

[13] R.H.W. ten Thije and R. Akkerman, Design of an experimental setup to measure tool-ply and ply-ply friction in thermoplastic laminates, Int. J. Mater. Form. 2:1 (2009) 197-200. https://doi.org/10.1007/s12289-009-0638-y

[14] P. Patel, T.R. Hull, R.W. McCabe, D. Flath, J. Grasmeder, and M. Percy, Mechanism of thermal decomposition of poly(ether ether ketone) (PEEK) from a review of decomposition studies, Polym. Degrad. Stab. 95:5 (2010) 709-718. https://doi.org/10.1016/j.polymdegradstab.2010.01.024

[15] C. Croshaw, L. Hamernik, L. Ghanbari, A. Browning, and J. Wiggins, Melt-state degradation mechanism of poly (ether ketone ketone): the role of branching on crystallization and rheological behavior, Polym. Degrad. Stab. 200 (2022) p. 109968. https://doi.org/10.1016/j.polymdegradstab.2022.109968

[16] M. Sonnenberg et al., Polymer-induced metal diffusion during plastic processing: a reason for deposit formation, J. Polym. Eng. 39:5 (2019) 472-480. https://doi.org/10.1515/polyeng-2018-0324

[17] C. Theile-Rasche, T. Meng, T. de los Arcos, and G. Grundmeier, "Analysis of polycarbonate degradation at melt/FeCr-alloy interfaces as a function of the alloy composition by means of combinatorial thin film chemistry," SN Appl. Sci. 5:10 (2023). https://doi.org/10.1007/s42452-023-05441-5

[18] J.M. Burkstrand, Formation of a copper-oxygen-polymer complex on polystyrene, Appl. Phys. Lett. 33:5 (1978) 387-389. https://doi.org/10.1063/1.90405

[19] J.M. Burkstrand, "Substrate effects on the electronic structure of metal overlayers--an XPS study of polymer-metal interfaces," Phys. Rev. B. 20:12 (1979) 4853-4858. https://doi.org/10.1103/PhysRevB.20.4853

[20] J.M. Burkstrand, "Metal-polymer interfaces: Adhesion and X-ray photoemission studies," J. Appl. Phys. 52:7 (1981) 4795-4800. https://doi.org/10.1063/1.329320

[21] M. Wiesing, T. de los Arcos, M. to Baben, H. Rueß, J. Schneider, and G. Grundmeier, Analysis of the inhibition of thermal degradation of molten polycarbonate at tool steel interfaces by thin TiAlN coatings, Polym. Degrad. Stab. 143 (2017) 196-206. https://doi.org/10.1016/j.polymdegradstab.2017.07.013

[22] G. Raos and B. Zappone, Polymer adhesion: Seeking new solutions for an old problem, Macromolecules 54:23 (2021) 10617-10644. https://doi.org/10.1021/acs.macromol.1c01182

[23] Data for this publication: https://doi.org/10.4121/f2a26463-a241-4698-bb3d-de31657d41d3
# An investigation into the relationship between hemisphere tooling parameters and thickness changes during the thermoforming of a thermoplastic composite

WHITE Kari D.¹,ᵃ* and SHERWOOD James A.¹,ᵇ

¹Department of Mechanical Engineering, University of Massachusetts, Lowell, One University Ave., 01854, Lowell, MA, USA

ᵃKari_White@uml.edu, ᵇJames_Sherwood@uml.edu

**Keywords:** Composite, Thermoforming, Simulation, Tooling, UHMWPE

**Abstract.** Thermoforming is an economical and fast means of transforming a flat sheet of thermoplastic composite material into a complex shape. The two-step process consists of a preform phase that transforms the laminates to near net shape ply stacks and of a subsequent consolidation step that employs pressure and heat to join the preforms into a final part. During the preform step, many aspects of the process must be regulated to minimize undesirable affects. A binder ring applies pressure to the material to induce in-plane tension, aiding in the initiation of shear deformation in the laminate. The material increases in thickness as it undergoes in-plane shear to conform to the preform shape. The degree of shear varies over the part, creating variations in thickness and subsequent challenges with uniform consolidation. While material attributes and processing parameters are critical in successful preforming, refining the tool geometry can also lead to significant changes in the preform outcome. Although many of these challenges have been addressed in a design-build-test approach, the goal of this research is to develop a virtual process that can guide design changes in the tooling to achieve a well consolidated part. A discrete mesoscopic modeling approach was implemented in LS-DYNA using thickness-change shell elements (ELFORM25) that can incorporate 3D material properties, including through-plane thickness changes. The current state of the simulation effectively incorporates material behavior, tool/ply friction and thickness changes that reflect variations in pressure. The simulation is used to investigate the significance of the punch/die gap size and the die fillet radius on the tooling geometry. The modeling approach was verified through comparison to experimental results with two types of tooling geometry. A parametric study was then performed to investigate tooling changes in relation to the number of layers, or thickness, of a unidirectional ultra-high molecular weight polyethylene (UHMWPE) material, DSM Dyneema® HB210. The outcomes of the result were evaluated by comparing wrinkle formation, shear deformation and thickness variations. The research shows that the radius of the die fillet has a significant impact on the uniformity of thickness in the preform. A combination of fillet radius and die affects can lead to more desirable outcomes in terms of wrinkle formation and thickness uniformity.

**Introduction**

Thermoforming is an attractive manufacturing process for high-volume, low-cost production of composites parts, and simulation is a valuable tool to guide the design of the processing parameters that can result in producing high-quality continuous fiber-reinforced composite parts. A typical thermoforming process consists of a preform step that transforms a set of laminates to a near-net-shape ply stack, and a subsequent consolidation phase that employs pressure and heat to join a set of the preforms into a final part. Friction plays a critical role in the preform phase of the manufacturing process.

<figure><img src="image_484.png" /></figure>

Content from this work may be used under the terms of the Creative Commons Attribution 3.0 license. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. Published under license by Materials Research Forum LLC.
# An investigation into the relationship between hemisphere tooling parameters and thickness changes during the thermoforming of a thermoplastic composite

WHITE Kari D.¹,ᵃ* and SHERWOOD James A.¹,ᵇ

¹Department of Mechanical Engineering, University of Massachusetts, Lowell, One University Ave., 01854, Lowell, MA, USA

ᵃKari_White@uml.edu, ᵇJames_Sherwood@uml.edu

**Keywords:** Composite, Thermoforming, Simulation, Tooling, UHMWPE

**Abstract.** Thermoforming is an economical and fast means of transforming a flat sheet of thermoplastic composite material into a complex shape. The two-step process consists of a preform phase that transforms the laminates to near net shape ply stacks and of a subsequent consolidation step that employs pressure and heat to join the preforms into a final part. During the preform step, many aspects of the process must be regulated to minimize undesirable affects. A binder ring applies pressure to the material to induce in-plane tension, aiding in the initiation of shear deformation in the laminate. The material increases in thickness as it undergoes in-plane shear to conform to the preform shape. The degree of shear varies over the part, creating variations in thickness and subsequent challenges with uniform consolidation. While material attributes and processing parameters are critical in successful preforming, refining the tool geometry can also lead to significant changes in the preform outcome. Although many of these challenges have been addressed in a design-build-test approach, the goal of this research is to develop a virtual process that can guide design changes in the tooling to achieve a well consolidated part. A discrete mesoscopic modeling approach was implemented in LS-DYNA using thickness-change shell elements (ELFORM25) that can incorporate 3D material properties, including through-plane thickness changes. The current state of the simulation effectively incorporates material behavior, tool/ply friction and thickness changes that reflect variations in pressure. The simulation is used to investigate the significance of the punch/die gap size and the die fillet radius on the tooling geometry. The modeling approach was verified through comparison to experimental results with two types of tooling geometry. A parametric study was then performed to investigate tooling changes in relation to the number of layers, or thickness, of a unidirectional ultra-high molecular weight polyethylene (UHMWPE) material, DSM Dyneema® HB210. The outcomes of the result were evaluated by comparing wrinkle formation, shear deformation and thickness variations. The research shows that the radius of the die fillet has a significant impact on the uniformity of thickness in the preform. A combination of fillet radius and die affects can lead to more desirable outcomes in terms of wrinkle formation and thickness uniformity.

**Introduction**

Thermoforming is an attractive manufacturing process for high-volume, low-cost production of composites parts, and simulation is a valuable tool to guide the design of the processing parameters that can result in producing high-quality continuous fiber-reinforced composite parts. A typical thermoforming process consists of a preform step that transforms a set of laminates to a near-net-shape ply stack, and a subsequent consolidation phase that employs pressure and heat to join a set of the preforms into a final part. Friction plays a critical role in the preform phase of the manufacturing process.

<figure><img src="image_495.png" /></figure>
Intra-ply shear is the dominant mode of deformation during the forming process, and various degrees of material shearing are what allows the material to conform across a compound curvature part. Adding complexity to the forming process outcome is the increase in thickness of the lamina due to conservation of volume during in-plane shear. Dangora et al. [1] documented the incompressibility phenomenon with micrographs taken of laminates of DSM Dyneema® HB80 sheared to 0°, 20°, and 60° showing the associated thickness increases corresponding to conservation of volume. Based on that research, variations in thickness can be linked to variations in shear. The implications of thickness changes during the process are two-fold: (1) thickness variations in the preforms and (2) effects on friction during the forming.

After the preform step, heat and pressure are applied in combination with matched dies to consolidate the preform(s) and to cure the part. Figure 1 depicts how thickness variations in individual preforms can lead to poor-quality consolidation and voids due to pressure variations. Therefore, these thickness changes must be considered in a forming simulation.

<figure><img src="image_485.png" /><figcaption>Figure 1. Variations in preform thickness leads to inconsistent / poor consolidation.</figcaption></figure>

The thickness changes due to shearing also affect material surfaces involved with frictional interactions. The friction between the binder ring and the material adds tension so that the material undergoes shear, even under the binder ring. Studies have shown that friction between two sliding surfaces can create unwanted outcomes such as distortions in the final part shape, residual stresses, and wrinkles [2-6]. Moreover, in the use of binder rings, this friction can help to reduce the formation of wrinkles with an appropriate pressure [1, 7-11]. Hemisphere preforming trials accomplished by Dangora et al. [1] showed the reliance on binder-ring pressure in regulating the size and appearance of wrinkles for a unidirectional cross-ply material system, DSM Dyneema® HB80.

However, tooling also contributes to the uniformity of preform thickness, the development of wrinkles, and the overall quality of the preformed part. The gap between the binder ring and the punch, the gap between the die and the punch, and the radii of both the binder ring and the die are parameters that can be altered to influence the preform results.

This paper presents a parametric study to study the influence of tooling geometry on the quality of a multi-layer thermoformed hemisphere. Simulation is used to examine the preform outcomes using a combination of binder and die geometries.

## Methodology

Experimental Hemisphere Preform. Hemisphere preform experiments were conducted using the two experimental setups shown in Figures 2 and 3. For both setups, the punch was attached to the crosshead of the Instron Universal Testing Machine, and load and displacement data were recorded. The hemisphere preform has a punch radius of 76.2 mm, and the laminate blanks used in the process are 381 mm by 381 mm. The crosshead descended at a rate of 6.35 mm/sec to replicate the conditions of the actual preform process.
Intra-ply shear is the dominant mode of deformation during the forming process, and various degrees of material shearing are what allows the material to conform across a compound curvature part. Adding complexity to the forming process outcome is the increase in thickness of the lamina due to conservation of volume during in-plane shear. Dangora et al. [1] documented the incompressibility phenomenon with micrographs taken of laminates of DSM Dyneema® HB80 sheared to 0°, 20°, and 60° showing the associated thickness increases corresponding to conservation of volume. Based on that research, variations in thickness can be linked to variations in shear. The implications of thickness changes during the process are two-fold: (1) thickness variations in the preforms and (2) effects on friction during the forming.

After the preform step, heat and pressure are applied in combination with matched dies to consolidate the preform(s) and to cure the part. Figure 1 depicts how thickness variations in individual preforms can lead to poor-quality consolidation and voids due to pressure variations. Therefore, these thickness changes must be considered in a forming simulation.

<figure><img src="image_496.png" /><figcaption>Figure 1. Variations in preform thickness leads to inconsistent / poor consolidation.</figcaption></figure>

The thickness changes due to shearing also affect material surfaces involved with frictional interactions. The friction between the binder ring and the material adds tension so that the material undergoes shear, even under the binder ring. Studies have shown that friction between two sliding surfaces can create unwanted outcomes such as distortions in the final part shape, residual stresses, and wrinkles [2-6]. Moreover, in the use of binder rings, this friction can help to reduce the formation of wrinkles with an appropriate pressure [1, 7-11]. Hemisphere preforming trials accomplished by Dangora et al. [1] showed the reliance on binder-ring pressure in regulating the size and appearance of wrinkles for a unidirectional cross-ply material system, DSM Dyneema® HB80.

However, tooling also contributes to the uniformity of preform thickness, the development of wrinkles, and the overall quality of the preformed part. The gap between the binder ring and the punch, the gap between the die and the punch, and the radii of both the binder ring and the die are parameters that can be altered to influence the preform results.

This paper presents a parametric study to study the influence of tooling geometry on the quality of a multi-layer thermoformed hemisphere. Simulation is used to examine the preform outcomes using a combination of binder and die geometries.

## Methodology

Experimental Hemisphere Preform. Hemisphere preform experiments were conducted using the two experimental setups shown in Figures 2 and 3. For both setups, the punch was attached to the crosshead of the Instron Universal Testing Machine, and load and displacement data were recorded. The hemisphere preform has a punch radius of 76.2 mm, and the laminate blanks used in the process are 381 mm by 381 mm. The crosshead descended at a rate of 6.35 mm/sec to replicate the conditions of the actual preform process.
<figure><img src="image_486.png" /><figcaption>Figure 2. Original experimental preform experimental setup.</figcaption></figure>

<figure><img src="image_487.png" /><figcaption>Figure 3. Redesigned experimental preform experimental setup.</figcaption></figure>

In the original setup, after the laminate sheets were in place, free weights were placed on top of the binder ring to develop pressure on the laminate. The entire setup was located inside an Instron oven to heat the material to the desired testing temperature. The maximum binder ring pressure achieved with the original setup is 5.75 kPa.

Finite element simulations showed that much higher binder pressures are required to mitigate wrinkle formation when considering multiple plies. The preform setup of Figure 2 was redesigned to accommodate more layers and to use a series of pneumatic cylinders to apply a greater force to the binder than was possible in the original design. The redesigned setup can apply pressures of up to a MPa to the laminate under the binder ring and includes a larger gap between the punch and the binder ring to accommodate up to ten layers of laminate. However, the redesigned test setup cannot fit in the Instron oven. Thus, the material is heated slightly above testing temperature in an infrared oven, and then quickly transferred to the new setup for forming.

**Modeling Approach.** The Sherwood Group [12] developed a discrete modeling approach that has been shown to successfully simulate the forming of a plain-weave hemisphere shape. Conventional elements available in commercial finite element software are used in the unit cell approach shown in Figure 4. Fiber directionality is described and tracked by beam elements that represent the tensile and flexural properties of the laminate. The shear load is carried by shell elements that capture the evolution of the in-plane shear stiffness as a function of the degree of shear. In the case of woven materials, the horizontal and vertical beams represent the warp and weft tows with of a bilinear modulus for the beam elements to capture wrinkling behavior [13]. In the case of a 0/90 cross-ply, the horizontal and vertical beams can be assigned tensile and flexural properties of the 0° and 90° layers.

<figure><img src="image_488.png" /><figcaption>Figure 4. Unit cell configuration for mesoscopic laminate material model.</figcaption></figure>

Plane-strain shell elements are widely used to model the in-plane shear behavior of laminate forming, but this class of shell elements does not capture the thickness changes caused by shearing, compression and/or stretching. The element formulations in LS-DYNA that can model through-plane strains were evaluated, and one formulation was chosen for further consideration. The LS-DYNA thickness-enhanced shell formulation (thick-thin shell), SHELL ELFORM25, is a general
<figure><img src="image_497.png" /><figcaption>Figure 2. Original experimental preform experimental setup.</figcaption></figure>

<figure><img src="image_498.png" /><figcaption>Figure 3. Redesigned experimental preform experimental setup.</figcaption></figure>

In the original setup, after the laminate sheets were in place, free weights were placed on top of the binder ring to develop pressure on the laminate. The entire setup was located inside an Instron oven to heat the material to the desired testing temperature. The maximum binder ring pressure achieved with the original setup is 5.75 kPa.

Finite element simulations showed that much higher binder pressures are required to mitigate wrinkle formation when considering multiple plies. The preform setup of Figure 2 was redesigned to accommodate more layers and to use a series of pneumatic cylinders to apply a greater force to the binder than was possible in the original design. The redesigned setup can apply pressures of up to a MPa to the laminate under the binder ring and includes a larger gap between the punch and the binder ring to accommodate up to ten layers of laminate. However, the redesigned test setup cannot fit in the Instron oven. Thus, the material is heated slightly above testing temperature in an infrared oven, and then quickly transferred to the new setup for forming.

Modeling Approach. The Sherwood Group [12] developed a discrete modeling approach that has been shown to successfully simulate the forming of a plain-weave hemisphere shape. Conventional elements available in commercial finite element software are used in the unit cell approach shown in Figure 4. Fiber directionality is described and tracked by beam elements that represent the tensile and flexural properties of the laminate. The shear load is carried by shell elements that capture the evolution of the in-plane shear stiffness as a function of the degree of shear. In the case of woven materials, the horizontal and vertical beams represent the warp and weft tows with of a bilinear modulus for the beam elements to capture wrinkling behavior [13]. In the case of a 0/90 cross-ply, the horizontal and vertical beams can be assigned tensile and flexural properties of the 0° and 90° layers.

<figure><img src="image_499.png" /><figcaption>Figure 4. Unit cell configuration for mesoscopic laminate material model.</figcaption></figure>

Plane-strain shell elements are widely used to model the in-plane shear behavior of laminate forming, but this class of shell elements does not capture the thickness changes caused by shearing, compression and/or stretching. The element formulations in LS-DYNA that can model through-plane strains were evaluated, and one formulation was chosen for further consideration. The LS-DYNA thickness-enhanced shell formulation (thick-thin shell), SHELL ELFORM25, is a general
shell element that includes an additional feature of linear strain through the thickness compared to a Hughes-Liu general shell element. Loading and contact of the surface are possible to allow for tooling and laminate plies to interact when contact algorithms are defined [14]. The thickness stretch of the element requires a 3D constitutive model, and the element is defined by a midsurface with only four nodes (Figure 5).

<figure><img src="image_489.png" /><figcaption>Figure 5. Thick-thin shell element.</figcaption></figure>

<figure><img src="image_490.png" /><figcaption>Figure 6. Hemisphere preform simulation in LS-DYNA.</figcaption></figure>

Evaluation of this element choice required that the user-defined (UMAT) discrete mesoscopic material model used in LS-DYNA consider 3D material behavior. The in-plane tensile behavior is still carried by beam elements that provide directionality to the elements carrying the shear load. The in-plane shear behavior as a function of shear angle remains the same as the plain-strain implementation. However, through-plane stress and strain are updated in the UMAT, whereas general shell elements do not have strain in that direction. Validation of the shear behavior application was done through comparison to shear-frame, also known as picture-frame, experimental data conducted with a constant shear rate of 0.1 sec⁻¹ [15]. The hemisphere preform simulation as shown in Figure 6 was modeled after the experimental setup, modified appropriately to represent the binder ring and die sizes. The friction coefficients for the material were calculated from data that was recorded using the UMass Lowell friction tester as redesigned by Campshure et al. [16-18].

The composite laminate system investigated in this research is DSM® Dyneema® HB210. This material system is a thermoplastic cross-ply containing four unidirectional layers oriented in a (0/90)₂ fiber configuration with each ply comprised of UHMWPE fibers and a thermoplastic polyurethane (TPU) based matrix [19]. The tensile, shear and bending properties at 100°C for Dyneema® HB210 are provided from previous testing [15, 20, 2] and shown in Table 1, where γ is defined as the shear strain of the composite lamina. The friction values are a result of averages between fiber orientation and resin condition, as the hemisphere is symmetrical with each condition represented equally.

<table><caption>Table 1. Material Properties for Hemisphere Preforming Simulations (MPa)</caption><thead><tr><td>Shear Stiffness as a Function of Shear Strain (MPa)</td><td>Tensile Modulus (MPa)</td><td>Compressive Modulus (MPa)</td><td>Tool/Ply Static</td><td>Tool/Ply Dynamic</td></tr></thead><tbody><tr><td>1723-5002|γ|⁵+5677|γ|⁴-3179|γ|³+909|γ|²-125|γ|+7</td><td>19300</td><td>5000</td><td>0.110</td><td>0.077</td></tr></tbody></table>

## Setup of Parametric Study on Tooling Geometry

Simulations of both experimental preforming setups were completed to validate the material properties and friction values used in the simulation. Measured punch force, shear angle distribution, draw-in profile, and visual wrinkle distribution were metrics used in the evaluation. All models showed good correlation with the experimental results, and it was concluded that the model was valid to move forward with the parametric study of the tooling.
shell element that includes an additional feature of linear strain through the thickness compared to a Hughes-Liu general shell element. Loading and contact of the surface are possible to allow for tooling and laminate plies to interact when contact algorithms are defined [14]. The thickness stretch of the element requires a 3D constitutive model, and the element is defined by a midsurface with only four nodes (Figure 5).

<figure><img src="image_500.png" /><figcaption>Figure 5. Thick-thin shell element.</figcaption></figure>

<figure><img src="image_501.png" /><figcaption>Figure 6. Hemisphere preform simulation in LS-DYNA.</figcaption></figure>

Evaluation of this element choice required that the user-defined (UMAT) discrete mesoscopic material model used in LS-DYNA consider 3D material behavior. The in-plane tensile behavior is still carried by beam elements that provide directionality to the elements carrying the shear load. The in-plane shear behavior as a function of shear angle remains the same as the plain-strain implementation. However, through-plane stress and strain are updated in the UMAT, whereas general shell elements do not have strain in that direction. Validation of the shear behavior application was done through comparison to shear-frame, also known as picture-frame, experimental data conducted with a constant shear rate of 0.1 sec⁻¹ [15]. The hemisphere preform simulation as shown in Figure 6 was modeled after the experimental setup, modified appropriately to represent the binder ring and die sizes. The friction coefficients for the material were calculated from data that was recorded using the UMass Lowell friction tester as redesigned by Campshure et al. [16-18].

The composite laminate system investigated in this research is DSM® Dyneema® HB210. This material system is a thermoplastic cross-ply containing four unidirectional layers oriented in a (0/90)₂ fiber configuration with each ply comprised of UHMWPE fibers and a thermoplastic polyurethane (TPU) based matrix [19]. The tensile, shear and bending properties at 100°C for Dyneema® HB210 are provided from previous testing [15, 20, 2] and shown in Table 1, where γ is defined as the shear strain of the composite lamina. The friction values are a result of averages between fiber orientation and resin condition, as the hemisphere is symmetrical with each condition represented equally.

<table><caption>Table 1. Material Properties for Hemisphere Preforming Simulations (MPa)</caption><thead><tr><td>Shear Stiffness as a Function of Shear Strain (MPa)</td><td>Tensile Modulus (MPa)</td><td>Compressive Modulus (MPa)</td><td>Tool/Ply Static</td><td>Tool/Ply Dynamic</td></tr></thead><tbody><tr><td>1723-5002|γ|⁵+5677|γ|⁴-3179|γ|³+909|γ|²-125|γ|+7</td><td>19300</td><td>5000</td><td>0.110</td><td>0.077</td></tr></tbody></table>

## Setup of Parametric Study on Tooling Geometry

Simulations of both experimental preforming setups were completed to validate the material properties and friction values used in the simulation. Measured punch force, shear angle distribution, draw-in profile, and visual wrinkle distribution were metrics used in the evaluation. All models showed good correlation with the experimental results, and it was concluded that the model was valid to move forward with the parametric study of the tooling.
Investigation of Tooling Parameters. Specific geometry attributes were identified to be relevant to the preforming outcomes. Figure 7 shows the upper gap as the distance between the inner radius of the binder ring and the outer radius of the punch. The lower gap indicated the distance between the inner radius of the die and the punch. The lower radius is a measure of the fillet, or curvature, of the die.

<figure><img src="image_491.png" /><figcaption>Figure 7. Tooling parameters analyzed in study.</figcaption></figure>

The experimental preform setups had the geometric values listed in Table 2, and the simulations were created to reflect these values. An initial parametric study was conducted that used combinations of the existing tooling to evaluate the sensitivity of the model to parameter changes. Based on the preliminary results, a set of tooling parameters was chosen to expand the range of binder gap (upper gap) and die fillet radius (lower radius). The range of values used in the study are listed in Tables 4 and 5. A binder pressure of 300 kPa was used to replicate the experimental pressure used with the four layers of laminate with the redesigned tooling. All simulations were performed with the properties of DSM® Dyneema® HB210 at a processing temperature of 100°C. Each simulation was a combination of a letter from Table 4 and a number from Table 5, e.g. “a1”.

<table><caption>Table 2. Geometric Properties of Experimental Preform Setup</caption><thead><tr><th>Tooling</th><th>Upper Gap<br>[mm]</th><th>Lower Gap<br>[mm]</th><th>Lower Radius<br>[mm]</th></tr></thead><tbody><tr><td>Original</td><td>8.26</td><td>3.18</td><td>12.7</td></tr><tr><td>Redesign</td><td>12.7</td><td>3.18</td><td>6.35</td></tr></tbody></table>

<table><caption>Table 3. Binder (upper) gap radius values</caption><thead><tr><th>Model<br>Designation</th><th>Upper Gap<br>[mm]</th></tr></thead><tbody><tr><td>a</td><td>7.62</td></tr><tr><td>b</td><td>8.26</td></tr><tr><td>c</td><td>10.2</td></tr><tr><td>d</td><td>12.7</td></tr><tr><td>e</td><td>15.2</td></tr></tbody></table>

<table><caption>Table 4. Die fillet (lower) radius values</caption><thead><tr><th>Model<br>Designation</th><th>Lower Radius<br>[mm]</th></tr></thead><tbody><tr><td>1</td><td>2.54</td></tr><tr><td>2</td><td>5.08</td></tr><tr><td>3</td><td>7.62</td></tr><tr><td>4</td><td>10.2</td></tr><tr><td>5</td><td>12.7</td></tr></tbody></table>
Investigation of Tooling Parameters. Specific geometry attributes were identified to be relevant to the preforming outcomes. Figure 7 shows the upper gap as the distance between the inner radius of the binder ring and the outer radius of the punch. The lower gap indicated the distance between the inner radius of the die and the punch. The lower radius is a measure of the fillet, or curvature, of the die.

<figure><img src="image_502.png" /><figcaption>Figure 7. Tooling parameters analyzed in study.</figcaption></figure>

The experimental preform setups had the geometric values listed in Table 2, and the simulations were created to reflect these values. An initial parametric study was conducted that used combinations of the existing tooling to evaluate the sensitivity of the model to parameter changes. Based on the preliminary results, a set of tooling parameters was chosen to expand the range of binder gap (upper gap) and die fillet radius (lower radius). The range of values used in the study are listed in Tables 4 and 5. A binder pressure of 300 kPa was used to replicate the experimental pressure used with the four layers of laminate with the redesigned tooling. All simulations were performed with the properties of DSM® Dyneema® HB210 at a processing temperature of 100°C. Each simulation was a combination of a letter from Table 4 and a number from Table 5, e.g. “a1”.

<table><caption>Table 2. Geometric Properties of Experimental Preform Setup</caption><thead><tr><th>Tooling</th><th>Upper Gap [mm]</th><th>Lower Gap [mm]</th><th>Lower Radius [mm]</th></tr></thead><tbody><tr><td>Original</td><td>8.26</td><td>3.18</td><td>12.7</td></tr><tr><td>Redesign</td><td>12.7</td><td>3.18</td><td>6.35</td></tr></tbody></table>

<table><caption>Table 3. Binder (upper) gap radius values</caption><thead><tr><th>Model Designation</th><th>Upper Gap [mm]</th></tr></thead><tbody><tr><td>a</td><td>7.62</td></tr><tr><td>b</td><td>8.26</td></tr><tr><td>c</td><td>10.2</td></tr><tr><td>d</td><td>12.7</td></tr><tr><td>e</td><td>15.2</td></tr></tbody></table>

<table><caption>Table 4. Die fillet (lower) radius values</caption><thead><tr><th>Model Designation</th><th>Lower Radius [mm]</th></tr></thead><tbody><tr><td>1</td><td>2.54</td></tr><tr><td>2</td><td>5.08</td></tr><tr><td>3</td><td>7.62</td></tr><tr><td>4</td><td>10.2</td></tr><tr><td>5</td><td>12.7</td></tr></tbody></table>
Three metrics were used in comparing the results from the study. Thickness distribution over the hemispherical portion of the preform gives guidance on how well the preform can be consolidated into a uniform part. As controlling thickness uniformity is a goal of this project, measures of thickness variation were used as a quantitative comparison. A general assessment of the quality of the preform includes the locations and amplitudes of wrinkles. Finally, the in-plane shear angle distribution contributes to both the thickness uniformity and an indication of how well the binder ring can engage appropriate shearing of the laminate into the hemispherical shape.

## Results and Discussion

The initial thickness of the material was 0.72 mm, and the final thicknesses of the elements over the hemisphere portion of the preform were normalized to this initial thickness to evaluate the relative change in thickness. Box and whisker plots were employed to visualize the distribution of element thickness over the part. Figure 8 shows normalized thickness results for all 25 tooling configurations that were analyzed. The box for each configuration represents the middle 50% of the data, while the lines extending to the top and bottom represent the max and min data values.

<figure><img src="image_492.png" /><figcaption>Figure 8. Normalized thickness variation for all tooling configurations.</figcaption></figure>

Thickness of the material over the preformed part is also tracked to indicate how well a set of preforms will be consolidated. The goal is to keep the variation in the thickness small. Therefore, the tighter (or narrower) the box and the smaller the range between the maximum and minimum the better. For example, configuration 5b in Figure 8 has a tight data spread while 5e has a large spread, and the thickness variation can be seen in the contour plots of Figure 9.

<figure><img src="image_493.png" /><figcaption>Figure 9. Thickness variations over hemisphere portion of preform for configurations (a) 5b and (b) 5e.</figcaption></figure>

The variations shown in Figure 8 demonstrate that tooling does affect the uniformity of thickness and choices can be made to improve that uniformity. Identifying the sensitivity of thickness variations to individual tooling components is critical. Figure 10 shows the 25 data sets parsed by the upper binder gap. It can be seen in this figure how the thickness uniformity changes as a function of changing the die fillet radius, with a subtle trend of increasing uniformity with
Three metrics were used in comparing the results from the study. Thickness distribution over the hemispherical portion of the preform gives guidance on how well the preform can be consolidated into a uniform part. As controlling thickness uniformity is a goal of this project, measures of thickness variation were used as a quantitative comparison. A general assessment of the quality of the preform includes the locations and amplitudes of wrinkles. Finally, the in-plane shear angle distribution contributes to both the thickness uniformity and an indication of how well the binder ring can engage appropriate shearing of the laminate into the hemispherical shape.

## Results and Discussion

The initial thickness of the material was 0.72 mm, and the final thicknesses of the elements over the hemisphere portion of the preform were normalized to this initial thickness to evaluate the relative change in thickness. Box and whisker plots were employed to visualize the distribution of element thickness over the part. Figure 8 shows normalized thickness results for all 25 tooling configurations that were analyzed. The box for each configuration represents the middle 50% of the data, while the lines extending to the top and bottom represent the max and min data values.

<figure><img src="image_503.png" /><figcaption>Figure 8. Normalized thickness variation for all tooling configurations.</figcaption></figure>

Thickness of the material over the preformed part is also tracked to indicate how well a set of preforms will be consolidated. The goal is to keep the variation in the thickness small. Therefore, the tighter (or narrower) the box and the smaller the range between the maximum and minimum the better. For example, configuration 5b in Figure 8 has a tight data spread while 5e has a large spread, and the thickness variation can be seen in the contour plots of Figure 9.

<figure><img src="image_504.png" /><figcaption>Figure 9. Thickness variations over hemisphere portion of preform for configurations (a) 5b and (b) 5e.</figcaption></figure>

The variations shown in Figure 8 demonstrate that tooling does affect the uniformity of thickness and choices can be made to improve that uniformity. Identifying the sensitivity of thickness variations to individual tooling components is critical. Figure 10 shows the 25 data sets parsed by the upper binder gap. It can be seen in this figure how the thickness uniformity changes as a function of changing the die fillet radius, with a subtle trend of increasing uniformity with
increasing fillet radius. However, not all the results fit the trend, indicating there may be competing effects with other tooling parameters or initial material thickness.

<figure><img src="image_494.png" /></figure>

<figure><img src="image_495.png" /><figcaption>(a) Upper binder gap 7.62 mm</figcaption></figure>

<figure><img src="image_496.png" /><figcaption>(b) Upper binder gap 8.26 mm</figcaption></figure>

<figure><img src="image_497.png" /><figcaption>(c) Upper binder gap 10.2 mm</figcaption></figure>

<figure><img src="image_498.png" /><figcaption>(d) Upper binder gap 12.7 mm</figcaption></figure>

<figure><img src="image_499.png" /><figcaption>(e) Upper binder gap 15.2 mm</figcaption></figure>

**Figure 10. Thickness variations as a function of increasing die fillet radius for five binder gaps**

(a-e).

Figure 11 shows the same 25 results when parsed by fillet radius. As the binder gap is increased, there is no recognizable trend in the thickness uniformity. These results imply that the fillet radius plays a larger role in controlling the uniformity of the preform thickness than the binder gap.

Based on the thickness uniformity results, the four configurations with the least variation in thickness were compared visually for part quality in terms of visible wrinkling. The preform surfaces of the four configurations are shown in Figure 12. Based on visible inspection, configurations 3d and 4a (Figure 12c) show the least amount of waviness and wrinkles extending up into the hemisphere from the base. Likewise, the amplitudes of the wrinkles appear to be less than the other configurations. Based on these results, specific combinations of binder gap and fillet radii would produce acceptable results. However, more research is necessary to quantify the amplitude and frequency of the wrinkles.

Figure 13 shows the shear angle contours for all four configurations. The shear results confirm that the material is being sheared adequately in the first two configurations, therefore the uniformity of material thickness is occurring because of the combination of the shear and tension
increasing fillet radius. However, not all the results fit the trend, indicating there may be competing effects with other tooling parameters or initial material thickness.

<figure><img src="image_505.png" /></figure>

<figure><img src="image_506.png" /><figcaption>(a) Upper binder gap 7.62 mm</figcaption></figure>

<figure><img src="image_507.png" /><figcaption>(b) Upper binder gap 8.26 mm</figcaption></figure>

<figure><img src="image_508.png" /><figcaption>(c) Upper binder gap 10.2 mm</figcaption></figure>

<figure><img src="image_509.png" /><figcaption>(d) Upper binder gap 12.7 mm</figcaption></figure>

<figure><img src="image_510.png" /><figcaption>(e) Upper binder gap 15.2mm</figcaption></figure>

**Figure 10. Thickness variations as a function of increasing die fillet radius for five binder gaps**

(a-e).

Figure 11 shows the same 25 results when parsed by fillet radius. As the binder gap is increased, there is no recognizable trend in the thickness uniformity. These results imply that the fillet radius plays a larger role in controlling the uniformity of the preform thickness than the binder gap.

Based on the thickness uniformity results, the four configurations with the least variation in thickness were compared visually for part quality in terms of visible wrinkling. The preform surfaces of the four configurations are shown in Figure 12. Based on visible inspection, configurations 3d and 4a (Figure 12c) show the least amount of waviness and wrinkles extending up into the hemisphere from the base. Likewise, the amplitudes of the wrinkles appear to be less than the other configurations. Based on these results, specific combinations of binder gap and fillet radii would produce acceptable results. However, more research is necessary to quantify the amplitude and frequency of the wrinkles.

Figure 13 shows the shear angle contours for all four configurations. The shear results confirm that the material is being sheared adequately in the first two configurations, therefore the uniformity of material thickness is occurring because of the combination of the shear and tension
effects. However, configurations with largest die fillet radius, the tension is not maintaining a level adequate for encouraging shearing of the material.

<figure><img src="image_500.png" /><figcaption>Figure 11. Thickness variations as a function of increasing binder gap for five fillet radii (a-e).</figcaption></figure>

<figure><img src="image_501.png" /></figure>
effects. However, configurations with largest die fillet radius, the tension is not maintaining a level adequate for encouraging shearing of the material.

<figure><img src="image_511.png" /><figcaption>Figure 11. Thickness variations as a function of increasing binder gap for five fillet radii (a-e).</figcaption></figure>

<figure><img src="image_512.png" /></figure>
<figure><img src="image_502.png" /><figcaption>Figure 13. Preform shear results for (a) 3d, (b)4a, (c) 5a and (d) 5b.</figcaption></figure>

Based on the preliminary studies, both the binder gap and die radius as well as combined affects contribute to the thickness uniformity and the wrinkle formation of the preform. Future work will strive to quantify the frequency and amplitude of the preform wrinkles for comparison.

## Summary

In this research hemisphere preforming simulations were performed on a unidirectional thermoplastic cross-ply material systems materials, DSM Dyneema® HB210. The simulation was successfully validated experimentally with two different tooling configurations. A parametric study was performed to evaluate the role of the binder gap (space between the punch and the binder) and the die radius (the curvature of the die that the material is pushed through). Results showed significant influence of lower radius in the uniformity of the thickness of the preform. The binder gap, in conjunction with the die fillet radius, plays a role in the wrinkle formation through appropriate shear in the material. Future studies will analyze the frequency and quality of wrinkling for quantitative comparison.

## Acknowledgement

This work was performed under Cooperative Agreement W911NF-18-2-0033 with the US Army Research Lab. The material was donated by DSM and LS-DYNA was donated by ANSYS.

## References

[1] Dangora, L.M., C.J. Mitchell, J. Sherwood, and J.C. Parker, Deep-Drawing Forming Trials on a Cross-Ply Thermoplastic Lamina for Helmet Preform Manufacture. Journal of Manufacturing Science and Engineering, 2017. 139(3): p. 031009. https://doi.org/10.1115/1.4034791

[2] Ersoy, N., T. Garstka, K. Potter, M.R. Wisnom, D. Porter, and G. Stringer, Modelling of the spring-in phenomenon in curved parts made of a thermosetting composite. Composites Part A: Applied Science and Manufacturing, 2010. 41(3): p. 410-418. https://doi.org/10.1016/j.compositesa.2009.11.008

[3] Ersoy, N., K. Potter, M.R. Wisnom, and M.J. Clegg, An experimental method to study the frictional processes during composites manufacturing. Composites Part A: Applied science and manufacturing, 2005. 36(11): p. 1536-1544. https://doi.org/10.1016/j.compositesa.2005.02.010

[4] Larberg, Y.R. and M. Åkermo, On the interplay friction of different generations of carbon/epoxy prepreg systems. Composites Part A: Applied Science and Manufacturing, 2011. 42(9): p. 1067-1074. https://doi.org/10.1016/j.compositesa.2011.04.010

[5] Lightfoot, J.S., M.R. Wisnom, and K. Potter, A new mechanism for the formation of ply wrinkles due to shear between plies. Composites Part A: Applied Science and Manufacturing, 2013. 49: p. 139-147. https://doi.org/10.1016/j.compositesa.2013.03.002

[6] Arafath, A.R.A., R. Vaziri, and A. Poursartip, Closed-form solution for process-induced stresses and deformation of a composite part cured on a solid tool: part I-flat geometries. Composites Part A: Applied Science and Manufacturing, 2008. 39(7): p. 1106-1117. https://doi.org/10.1016/j.compositesa.2008.04.009
<figure><img src="image_513.png" /><figcaption>Figure 13. Preform shear results for (a) 3d, (b)4a, (c) 5a and (d) 5b.</figcaption></figure>

Based on the preliminary studies, both the binder gap and die radius as well as combined affects contribute to the thickness uniformity and the wrinkle formation of the preform. Future work will strive to quantify the frequency and amplitude of the preform wrinkles for comparison.

**Summary**

In this research hemisphere preforming simulations were performed on a unidirectional thermoplastic cross-ply material systems materials, DSM Dyneema® HB210. The simulation was successfully validated experimentally with two different tooling configurations. A parametric study was performed to evaluate the role of the binder gap (space between the punch and the binder) and the die radius (the curvature of the die that the material is pushed through). Results showed significant influence of lower radius in the uniformity of the thickness of the preform. The binder gap, in conjunction with the die fillet radius, plays a role in the wrinkle formation through appropriate shear in the material. Future studies will analyze the frequency and quality of wrinkling for quantitative comparison.

**Acknowledgement**

This work was performed under Cooperative Agreement W911NF-18-2-0033 with the US Army Research Lab. The material was donated by DSM and LS-DYNA was donated by ANSYS.

**References**

[1] Dangora, L.M., C.J. Mitchell, J. Sherwood, and J.C. Parker, Deep-Drawing Forming Trials on a Cross-Ply Thermoplastic Lamina for Helmet Preform Manufacture. Journal of Manufacturing Science and Engineering, 2017. 139(3): p. 031009. https://doi.org/10.1115/1.4034791

[2] Ersoy, N., T. Garstka, K. Potter, M.R. Wisnom, D. Porter, and G. Stringer, Modelling of the spring-in phenomenon in curved parts made of a thermosetting composite. Composites Part A: Applied Science and Manufacturing, 2010. 41(3): p. 410-418. https://doi.org/10.1016/j.compositesa.2009.11.008

[3] Ersoy, N., K. Potter, M.R. Wisnom, and M.J. Clegg, An experimental method to study the frictional processes during composites manufacturing. Composites Part A: Applied science and manufacturing, 2005. 36(11): p. 1536-1544. https://doi.org/10.1016/j.compositesa.2005.02.010

[4] Larberg, Y.R. and M. Åkermo, On the interply friction of different generations of carbon/epoxy prepreg systems. Composites Part A: Applied Science and Manufacturing, 2011. 42(9): p. 1067-1074. https://doi.org/10.1016/j.compositesa.2011.04.010

[5] Lightfoot, J.S., M.R. Wisnom, and K. Potter, A new mechanism for the formation of ply wrinkles due to shear between plies. Composites Part A: Applied Science and Manufacturing, 2013. 49: p. 139-147. https://doi.org/10.1016/j.compositesa.2013.03.002

[6] Arafath, A.R.A., R. Vaziri, and A. Poursartip, Closed-form solution for process-induced stresses and deformation of a composite part cured on a solid tool: part I-flat geometries. Composites Part A: Applied Science and Manufacturing, 2008. 39(7): p. 1106-1117. https://doi.org/10.1016/j.compositesa.2008.04.009
[7] Breuer, U., M. Neitzel, V. Ketzer, and R. Reinicke, Deep drawing of fabric-reinforced thermoplastics: Wrinkle formation and their reduction. Polymer composites, 1996. 17(4): p. 643-647. https://doi.org/10.1002/pc.10655

[8] Lee, J.S., S.J. Hong, W.-R. Yu, and T.J. Kang, The effect of blank holder force on the stamp forming behavior of non-crimp fabric with a chain stitch. Composites science and technology, 2007. 67(3-4): p. 357-366. https://doi.org/10.1016/j.compscitech.2006.09.009

[9] Lin, H., J. Wang, A. Long, M. Clifford, and P. Harrison, Predictive modelling for optimization of textile composite forming. Composites Science and Technology, 2007. 67(15-16): p. 3242-3252. https://doi.org/10.1016/j.compscitech.2007.03.040

[10] Obermeyer, E. and S.A. Majlessi, A review of recent advances in the application of blank-holder force towards improving the forming limits of sheet metal parts. Journal of Materials Processing Technology, 1998. 75(1-3): p. 222-234. https://doi.org/10.1016/S0924-0136(97)00368-3

[11] Akkerman, R., R. Ten Thije, U. Sachs, and M. De Rooij. Friction in textile thermoplastic composites forming. in Proceedings of the 10th international conference on textile composites-TEXCOMP. 2010.

[12] Jauffres, D., J. Sherwood, C. Morris, and J. Chen, Discrete mesoscopic modeling for the simulation of woven-fabric reinforcement forming. Vol. 3. 2010. 1205-1216. https://doi.org/10.1007/s12289-009-0646-y

[13] Dangora, L.M., C.J. Mitchell, and J.A. Sherwood, Predictive model for the detection of out-of-plane defects formed during textile-composite manufacture. Composites Part A: Applied Science and Manufacturing, 2015. 78: p. 102-112. https://doi.org/10.1016/j.compositesa.2015.07.011

[14] Haufe, A., K. Schweizerhof, and P. Dubois, Properties & Limits: Review of Shell Element Formulations. 2013: https://www.researchgate.net/publication/312053480_Properties_Limits_Review_of_Shell_Elem ent_Formulations.

[15] White, K.D., C. Krogh, and J.A. Sherwood. Investigation of shear characterization of a UHMWPE unidirectional cross-ply for finite element simulation of composite processing. in AIP Conference Proceedings. 2019. AIP Publishing LLC. https://doi.org/10.1063/1.5112522

[16] Fetfatsidis, K.A., L.M. Gamache, J.L. Gorczyca, J.A. Sherwood, D. Jauffrès, and J. Chen, Design of an apparatus for measuring tool/fabric and fabric/fabric friction of woven-fabric composites during the therostamping process. International Journal of Material Forming, 2013. 6(1): p. 1-11. https://doi.org/10.1007/s12289-011-1058-3

[17] Campshure, B., K.D. White, and J.A. Sherwood, Friction Characterization of UHMWPE Cross-Ply Composite Sheets for Thermoforming Processes. 2021. https://doi.org/10.25518/esaform21.759

[18] White, K.D., B. Campshure, and J.A. Sherwood, Effects of Thickness Changes and Friction during the Thermoforming of Composite Sheets. 2021. https://doi.org/10.25518/esaform21.506

[19] DSM Dyneema® Industries. 2018 [cited 2017 June]; Available from: http://www.dsm.com/products/dyneema/en_US/home.html.

[20] White, K., M. Yeager, J. Sherwood, T. Bogetti, and J. Cline, Material Characterization and Finite Element Modeling for the Forming of Highly Oriented UHMWPE Thin Film and Unidirectional Cross-ply Composites, in ASC 33rd Technical Conference. 2018: Seattle, WA. https://doi.org/10.12783/asc33/25944

[21] White, K.D., L. Dangora, and J.A. Sherwood, Temperature Dependent Flexural Rigidities and Thickness Investigation, in ASC 31st Technical Conference. 2017: Williamsburg, VA.
[7] Breuer, U., M. Neitzel, V. Ketzer, and R. Reinicke, Deep drawing of fabric-reinforced thermoplastics: Wrinkle formation and their reduction. Polymer composites, 1996. 17(4): p. 643-647. https://doi.org/10.1002/pc.10655

[8] Lee, J.S., S.J. Hong, W.-R. Yu, and T.J. Kang, The effect of blank holder force on the stamp forming behavior of non-crimp fabric with a chain stitch. Composites science and technology, 2007. 67(3-4): p. 357-366. https://doi.org/10.1016/j.compscitech.2006.09.009

[9] Lin, H., J. Wang, A. Long, M. Clifford, and P. Harrison, Predictive modelling for optimization of textile composite forming. Composites Science and Technology, 2007. 67(15-16): p. 3242-3252. https://doi.org/10.1016/j.compscitech.2007.03.040

[10] Obermeyer, E. and S.A. Majlessi, A review of recent advances in the application of blank-holder force towards improving the forming limits of sheet metal parts. Journal of Materials Processing Technology, 1998. 75(1-3): p. 222-234. https://doi.org/10.1016/S0924-0136(97)00368-3

[11] Akkerman, R., R. Ten Thije, U. Sachs, and M. De Rooij. Friction in textile thermoplastic composites forming. in Proceedings of the 10th international conference on textile composites-TEXCOMP. 2010.

[12] Jauffres, D., J. Sherwood, C. Morris, and J. Chen, Discrete mesoscopic modeling for the simulation of woven-fabric reinforcement forming. Vol. 3. 2010. 1205-1216. https://doi.org/10.1007/s12289-009-0646-y

[13] Dangora, L.M., C.J. Mitchell, and J.A. Sherwood, Predictive model for the detection of out-of-plane defects formed during textile-composite manufacture. Composites Part A: Applied Science and Manufacturing, 2015. 78: p. 102-112. https://doi.org/10.1016/j.compositesa.2015.07.011

[14] Haufe, A., K. Schweizerhof, and P. Dubois, Properties & Limits: Review of Shell Element Formulations. 2013: https://www.researchgate.net/publication/312053480_Properties_Limits_Review_of_Shell_Elem ent_Formulations.

[15] White, K.D., C. Krogh, and J.A. Sherwood. Investigation of shear characterization of a UHMWPE unidirectional cross-ply for finite element simulation of composite processing. in AIP Conference Proceedings. 2019. AIP Publishing LLC. https://doi.org/10.1063/1.5112522

[16] Fetfatsidis, K.A., L.M. Gamache, J.L. Gorczyca, J.A. Sherwood, D. Jauffrès, and J. Chen, Design of an apparatus for measuring tool/fabric and fabric/fabric friction of woven-fabric composites during the therostamping process. International Journal of Material Forming, 2013. 6(1): p. 1-11. https://doi.org/10.1007/s12289-011-1058-3

[17] Campshure, B., K.D. White, and J.A. Sherwood, Friction Characterization of UHMWPE Cross-Ply Composite Sheets for Thermoforming Processes. 2021. https://doi.org/10.25518/esaform21.759

[18] White, K.D., B. Campshure, and J.A. Sherwood, Effects of Thickness Changes and Friction during the Thermoforming of Composite Sheets. 2021. https://doi.org/10.25518/esaform21.506

[19] DSM Dyneema® Industries. 2018 [cited 2017 June]; Available from: http://www.dsm.com/products/dyneema/en_US/home.html.

[20] White, K., M. Yeager, J. Sherwood, T. Bogetti, and J. Cline, Material Characterization and Finite Element Modeling for the Forming of Highly Oriented UHMWPE Thin Film and Unidirectional Cross-ply Composites, in ASC 33rd Technical Conference. 2018: Seattle, WA. https://doi.org/10.12783/asc33/25944

[21] White, K.D., L. Dangora, and J.A. Sherwood, Temperature Dependent Flexural Rigidities and Thickness Investigation, in ASC 31st Technical Conference. 2017: Williamsburg, VA.
# Effects of UD and twill reinforcements in hybrid sheet molding compound laminates

RAIMONDI Luca¹,ª*, BRUGO Tommaso Maria¹,ᵇ,
ZUCCHELLI Andrea¹,ᶜ and DONATI Lorenzo¹,ᵈ

¹Viale Risorgimento 2, Bologna, Italy

ªluca.raimondi@unibo.it, ᵇtommasomaria.brugo@unibo.it, ᶜa.zucchelli@unibo.it,

ᵈI.donati@unibo.it

**Keywords:** Carbon Fiber, Compression Molding, Sheet Molding Compound, Hybrid Composites

**Abstract.** This work proposes a novel manufacturing solution to locally reinforce a Sheet Molding Compound (SMC) with continuous fiber prepreg in a single-step compression molding operation. An adapted SMC process allowing for the manufacturing of continuous carbon fiber-reinforced SMC was successfully implemented, and a layer of fabric was added to mitigate the waviness induced during processing. Plates with different stacking sequences were manufactured and used for mechanical characterization to understand the effect of hybridization over tensile and flexural properties. It emerges that hybridization led to a 137% increase in tensile strength and up to 147% rise in flexural modulus compared to the standard SMC baseline. Additionally, the hybridization process helped reduce the variability in elastic and strength properties.

## Introduction

Driven by the need to reduce operating costs and increase performance and safety, the transport industry is looking for advanced, high-specific strength materials to meet cost, quality, and volume targets. Discontinued fiber-reinforced composite materials are appreciated for their high processability, high stiffness [1], notch insensitivity [2], and ability to withstand harsh environmental conditions [3]. However, they exhibit relatively low strength due to their limited reinforcement length. Unidirectional and continuous fiber-reinforced composites are highly regarded for their impressive specific stiffness and strength as well as for their capability of energy absorption in the event of a crash [4,5], but have limited design flexibility due to poor formability. The concept of integrating continuous reinforcements with discontinuous fiber reinforced in the form of Sheet Molding Compound (SMC) presents a promising solution to overcome the inherent limitations of both materials. While the concept of blending discontinuous SMC with continuous fibrous reinforcements originated in the 1970s [6,7], to the best of the authors' knowledge, comprehensive studies specifically focused on understanding hybrid SMC materials have been conducted for approximately the past five years. A potential limitation dictated by the technology arises from the opposite need to establish appropriate flow in the short-fiber layers and to limit the mobility of the long-fiber composite layers to prevent misalignment phenomena during molding. Strategies to mitigate these defects include staging the long fiber reinforced prepreg resin before molding [9] or using dual-stage resin systems [8,9]. Although these strategies have proven effective against defect formation, they require prior processing in an oven or to alter the temperature of the tool during forming, which slows down the production time of the manufactured part. The possibility of manufacturing hybrid laminates in a single-step compression molding operation is investigated in this work. An adapted SMC process allowing for the manufacturing of continuous carbon fiber reinforced SMC was successfully implemented, and a layer of fabric was added to mitigate the waviness induced during processing. Plates with different stacking sequences were manufactured and used for mechanical characterization to understand the effect of

<figure><img src="image_503.png" /></figure>
# Effects of UD and twill reinforcements in hybrid sheet molding compound laminates

RAIMONDI Luca¹,ª*, BRUGO Tommaso Maria¹,ᵇ,
ZUCCHELLI Andrea¹,ᶜ and DONATI Lorenzo¹,ᵈ

¹Viale Risorgimento 2, Bologna, Italy

ªluca.raimondi@unibo.it, ᵇtommasomaria.brugo@unibo.it, ᶜa.zucchelli@unibo.it,

ᵈI.donati@unibo.it

**Keywords:** Carbon Fiber, Compression Molding, Sheet Molding Compound, Hybrid Composites

**Abstract.** This work proposes a novel manufacturing solution to locally reinforce a Sheet Molding Compound (SMC) with continuous fiber prepreg in a single-step compression molding operation. An adapted SMC process allowing for the manufacturing of continuous carbon fiber-reinforced SMC was successfully implemented, and a layer of fabric was added to mitigate the waviness induced during processing. Plates with different stacking sequences were manufactured and used for mechanical characterization to understand the effect of hybridization over tensile and flexural properties. It emerges that hybridization led to a 137% increase in tensile strength and up to 147% rise in flexural modulus compared to the standard SMC baseline. Additionally, the hybridization process helped reduce the variability in elastic and strength properties.

## Introduction

Driven by the need to reduce operating costs and increase performance and safety, the transport industry is looking for advanced, high-specific strength materials to meet cost, quality, and volume targets. Discontinued fiber-reinforced composite materials are appreciated for their high processability, high stiffness [1], notch insensitivity [2], and ability to withstand harsh environmental conditions [3]. However, they exhibit relatively low strength due to their limited reinforcement length. Unidirectional and continuous fiber-reinforced composites are highly regarded for their impressive specific stiffness and strength as well as for their capability of energy absorption in the event of a crash [4,5], but have limited design flexibility due to poor formability. The concept of integrating continuous reinforcements with discontinuous fiber reinforced in the form of Sheet Molding Compound (SMC) presents a promising solution to overcome the inherent limitations of both materials. While the concept of blending discontinuous SMC with continuous fibrous reinforcements originated in the 1970s [6,7], to the best of the authors' knowledge, comprehensive studies specifically focused on understanding hybrid SMC materials have been conducted for approximately the past five years. A potential limitation dictated by the technology arises from the opposite need to establish appropriate flow in the short-fiber layers and to limit the mobility of the long-fiber composite layers to prevent misalignment phenomena during molding. Strategies to mitigate these defects include staging the long fiber reinforced prepreg resin before molding [9] or using dual-stage resin systems [8,9]. Although these strategies have proven effective against defect formation, they require prior processing in an oven or to alter the temperature of the tool during forming, which slows down the production time of the manufactured part. The possibility of manufacturing hybrid laminates in a single-step compression molding operation is investigated in this work. An adapted SMC process allowing for the manufacturing of continuous carbon fiber reinforced SMC was successfully implemented, and a layer of fabric was added to mitigate the waviness induced during processing. Plates with different stacking sequences were manufactured and used for mechanical characterization to understand the effect of

<figure><img src="image_514.png" /></figure>
hybridization over tensile and flexural properties. Before mechanical testing, fiber misalignment of the endless fiber reinforcements was characterized by the Full Field Nodal Method (FFNM) [10] to measure the quality of the produced plates.

## Materials and Methods

*Materials, plate manufacturing, and sample extraction.* In this study, three materials were used: a sheet molding compound (SMC) prepreg, a unidirectional prepreg (UD), and a twill (TW) prepreg. The SMC used was HexMC/C/2000/M77, supplied by Hexcel (Duxford, U.K.). This SMC comprises randomly oriented 25 mm long carbon fiber strands impregnated with M77 epoxy resin. M77 is a fast-curing epoxy specifically designed for compression molding. The UD used was Toray prepreg P 384_S, supplied by Torayca (Tokyo, Japan). It is made of high-strength carbon fibers and impregnated with 40% #2300 resin, both manufactured by Toray Industries. The TW prepreg used was a 2x2 fabric SC160/RC200P supplied by Gurit (Volpiano, IT). This TW is made of standard modulus carbon fibers manufactured by Toray Industries and impregnated with a 40% GURIT SC160 epoxy resin. Different charge configurations for misalignment analysis and mechanical characterization of the three materials and their combination were considered. A mold coverage of 100% was adopted to manufacture the continuous fiber plates to prevent material flow during the process. The prepreg was cut into 320 [mm] x 240 [mm] rectangular preforms to manufacture the SMC plaques, achieving 80% tool coverage and allowing a one-dimensional flow during compression molding. UD and TW patches were first laminated with 100% areal coverage to produce the continuous reinforcement. The protocol described in [10,11] was adopted to measure the process induced waviness of external plies. An uniform 35 [mm] x 35 [mm] grid was drawn on the surface of the wet prepreg by using a 3 mm wide Berner paint marker. Images of gridded uncured plates were acquired into a Kyocera TASKalfa 3501i Multifunction Laser Printer at a resolution of 600 × 600 dpi and stored for subsequent processing. Afterwards, the charge was completed by stacking an 80% areal coverage SMC laminate and then compression molded in a one-step compression molding operation as schematically represented in Figure 1.

<figure><img src="image_504.png" /><figcaption>Figure 1: Compression molding of hybrid plates.</figcaption></figure>

An overview of the staking sequence and designation used for each plaque's charge is provided in Table 1. For example, UD2/100%/SMC2/80% consists of two plies of UD having a tool coverage area of 100% stacked onto two plies of SMC with an areal coverage of 80%. The compression molding was performed using a specialized tool measuring 320 [mm] x 320 [mm], designed to produce plates with arbitrary thicknesses. The tool was installed in a downstroke AEM3 press, and its temperature was maintained at a constant 130°C, which was verified by two thermocouples. Once cured, the plaques were extracted, cooled in free air until they reached room temperature, and scanned using a resolution of 600 × 600 dpi. The size of the plaques allows for the extraction of at least 5 samples for each mechanical testing configuration. The specimens were cut using abrasive water jet cutting and stored at room temperature with medium relative humidity for seven days before testing.
hybridization over tensile and flexural properties. Before mechanical testing, fiber misalignment of the endless fiber reinforcements was characterized by the Full Field Nodal Method (FFNM) [10] to measure the quality of the produced plates.

## Materials and Methods

*Materials, plate manufacturing, and sample extraction.* In this study, three materials were used: a sheet molding compound (SMC) prepreg, a unidirectional prepreg (UD), and a twill (TW) prepreg. The SMC used was HexMC/C/2000/M77, supplied by Hexcel (Duxford, U.K.). This SMC comprises randomly oriented 25 mm long carbon fiber strands impregnated with M77 epoxy resin. M77 is a fast-curing epoxy specifically designed for compression molding. The UD used was Toray prepreg P 384_S, supplied by Torayca (Tokyo, Japan). It is made of high-strength carbon fibers and impregnated with 40% #2300 resin, both manufactured by Toray Industries. The TW prepreg used was a 2x2 fabric SC160/RC200P supplied by Gurit (Volpiano, IT). This TW is made of standard modulus carbon fibers manufactured by Toray Industries and impregnated with a 40% GURIT SC160 epoxy resin. Different charge configurations for misalignment analysis and mechanical characterization of the three materials and their combination were considered. A mold coverage of 100% was adopted to manufacture the continuous fiber plates to prevent material flow during the process. The prepreg was cut into 320 [mm] x 240 [mm] rectangular preforms to manufacture the SMC plaques, achieving 80% tool coverage and allowing a one-dimensional flow during compression molding. UD and TW patches were first laminated with 100% areal coverage to produce the continuous reinforcement. The protocol described in [10,11] was adopted to measure the process induced waviness of external plies. An uniform 35 [mm] x 35 [mm] grid was drawn on the surface of the wet prepreg by using a 3 mm wide Berner paint marker. Images of gridded uncured plates were acquired into a Kyocera TASKalfa 3501i Multifunction Laser Printer at a resolution of 600 × 600 dpi and stored for subsequent processing. Afterwards, the charge was completed by stacking an 80% areal coverage SMC laminate and then compression molded in a one-step compression molding operation as schematically represented in Figure 1.

<figure><img src="image_515.png" /><figcaption>Figure 1: Compression molding of hybrid plates.</figcaption></figure>

An overview of the staking sequence and designation used for each plaque's charge is provided in Table 1. For example, UD2/100%/SMC2/80% consists of two plies of UD having a tool coverage area of 100% stacked onto two plies of SMC with an areal coverage of 80%. The compression molding was performed using a specialized tool measuring 320 [mm] x 320 [mm], designed to produce plates with arbitrary thicknesses. The tool was installed in a downstroke AEM3 press, and its temperature was maintained at a constant 130°C, which was verified by two thermocouples. Once cured, the plaques were extracted, cooled in free air until they reached room temperature, and scanned using at a resolution of 600 × 600 dpi. The size of the plaques allows for the extraction of at least 5 samples for each mechanical testing configuration. The specimens were cut using abrasive water jet cutting and stored at room temperature with medium relative humidity for seven days before testing.
<table><caption>Table 1: Plaque stacking and molding details for different material combinations.</caption><thead><tr><th>Study #.<br>Plaque #</th><th>Designation</th><th># UD<br>plies</th><th># TW<br>plies</th><th># SMC<br>plies</th><th>Charge coverage</th></tr></thead><tbody><tr><td>1.1</td><td>SMC3/80%</td><td></td><td></td><td>3</td><td>80%</td></tr><tr><td>1.2</td><td>SMC8/80%</td><td></td><td></td><td>8</td><td>80%</td></tr><tr><td>2.1</td><td>UD2/100%</td><td>2</td><td></td><td></td><td>100%</td></tr><tr><td>2.2</td><td>UD14/100%</td><td>14</td><td></td><td></td><td>100%</td></tr><tr><td>3.1</td><td>TW6/100%</td><td></td><td>6</td><td></td><td>100%</td></tr><tr><td>3.2</td><td>TW10/100%</td><td></td><td>10</td><td></td><td>100%</td></tr><tr><td>4.1</td><td>UD2/100%/TW1/100%/SMC2/80%</td><td>2</td><td>1</td><td>2</td><td>100%/100%/80%</td></tr></tbody></table>

In order to avoid any potential damage to the outer layers of continuous carbon fiber-reinforced material during testing, end tabs were used on tensile samples. To prepare for adhesive bonding, both sides of the samples were conditioned with 180-grit sandpaper and cleaned with isopropyl alcohol. Tapered SMC end tabs, which were 2.5 [mm] thick, 50 [mm] long, and 25 [mm] wide, were then bonded onto the specimens.

*Mechanical testing.* The specimens for tensile testing were designed according to the standard test method ASTM D3039/D3039M and featured a length of 250 [mm] and a width of 25 [mm]. Tests were conducted at a constant cross-head rate of 2.5 [mm/min] using a servo-hydraulic INSTRON 8033 universal testing machine equipped with a 250 [kN] load cell under displacement-controlled conditions. The sample's surfaces were monitored by the two cameras of a commercial 3D Digital Image Correlation (DIC) system (Q-400, Dantec Dynamics, Skovlunde, Denmark) equipped with 17-mm lenses (Xenoplan, Schneider-Kreuznach, Bad Kreuznach, Germany) to allow for reliable measurement of surface deformation as suggested by [12]. The specimen surface was coated with a water-based paint (black matt, Chreon, Lechler, Como, Italy). For the speckle dots, white paint was used (white matt, Chreon, Lechler). The paints were deposited on the specimens by means of an airbrush spray gun (AZ3 HTE 2, Antes Iwata, Torino, Italy) with a nozzle of 1.8 [mm], a round jet, and a spray distance of about 300 [mm]. The Istra-4D software from Dantec Dynamics was used to analyze images, determine displacements, and strain distributions. The facet size was set to 25 pixels, and the facet overlap to 9 pixels. A local regression displacement smoothing filter of 25 x 25 facets was applied to balance accuracy and spatial resolution [10]. The specimens for flexural testing were designed according to the standard test method ASTM D790-17. To minimize through-thickness shear effects, a span-to-thickness ratio greater than 16:1 was adopted. Specimens were loaded to failure at a constant cross-head rate of 1 [mm/min] using a 2 [kN] load cell.

## Result and discussion

Results from the misalignment analysis by FFNM on representative plates is given in Figure 2. Due to the intense processing pressure, both UD2/100% and UD14/100% plaques exhibited an in-plane fiber waviness in the range [-5° 5°] (Figure 2a), demonstrating that this material was able to flow even at 100% areal coverage. Apparent in-plane waviness was much less pronounced in TW6/100% and TW10/100% plaques, being consistently in the range of [0° 2°] as reported in Figure 2b. In-plane fiber waviness of the outer layer of UD2/100%/TW1/100%/SMC2/80% hybrid plaques was consistently in the range of [-3° 2°], demonstrating the effectiveness of TW fabric in preventing lateral motion of UD fibers during processing.
<table><caption>Table 1: Plaque stacking and molding details for different material combinations.</caption><thead><tr><th>Study #.<br>Plaque #</th><th>Designation</th><th># UD<br>plies</th><th># TW<br>plies</th><th># SMC<br>plies</th><th>Charge coverage</th></tr></thead><tbody><tr><td>1.1</td><td>SMC3/80%</td><td></td><td></td><td>3</td><td>80%</td></tr><tr><td>1.2</td><td>SMC8/80%</td><td></td><td></td><td>8</td><td>80%</td></tr><tr><td>2.1</td><td>UD2/100%</td><td>2</td><td></td><td></td><td>100%</td></tr><tr><td>2.2</td><td>UD14/100%</td><td>14</td><td></td><td></td><td>100%</td></tr><tr><td>3.1</td><td>TW6/100%</td><td></td><td>6</td><td></td><td>100%</td></tr><tr><td>3.2</td><td>TW10/100%</td><td></td><td>10</td><td></td><td>100%</td></tr><tr><td>4.1</td><td>UD2/100%/TW1/100%/SMC2/80%</td><td>2</td><td>1</td><td>2</td><td>100%/100%/80%</td></tr></tbody></table>

In order to avoid any potential damage to the outer layers of continuous carbon fiber-reinforced material during testing, end tabs were used on tensile samples. To prepare for adhesive bonding, both sides of the samples were conditioned with 180-grit sandpaper and cleaned with isopropyl alcohol. Tapered SMC end tabs, which were 2.5 [mm] thick, 50 [mm] long, and 25 [mm] wide, were then bonded onto the specimens.

*Mechanical testing.* The specimens for tensile testing were designed according to the standard test method ASTM D3039/D3039M and featured a length of 250 [mm] and a width of 25 [mm]. Tests were conducted at a constant cross-head rate of 2.5 [mm/min] using a servo-hydraulic INSTRON 8033 universal testing machine equipped with a 250 [kN] load cell under displacement-controlled conditions. The sample's surfaces were monitored by the two cameras of a commercial 3D Digital Image Correlation (DIC) system (Q-400, Dantec Dynamics, Skovlunde, Denmark) equipped with 17-mm lenses (Xenoplan, Schneider-Kreuznach, Bad Kreuznach, Germany) to allow for reliable measurement of surface deformation as suggested by [12]. The specimen surface was coated with a water-based paint (black matt, Chreon, Lechler, Como, Italy). For the speckle dots, white paint was used (white matt, Chreon, Lechler). The paints were deposited on the specimens by means of an airbrush spray gun (AZ3 HTE 2, Antes Iwata, Torino, Italy) with a nozzle of 1.8 [mm], a round jet, and a spray distance of about 300 [mm]. The Istra-4D software from Dantec Dynamics was used to analyze images, determine displacements, and strain distributions. The facet size was set to 25 pixels, and the facet overlap to 9 pixels. A local regression displacement smoothing filter of 25 x 25 facets was applied to balance accuracy and spatial resolution [10]. The specimens for flexural testing were designed according to the standard test method ASTM D790-17. To minimize through-thickness shear effects, a span-to-thickness ratio greater than 16:1 was adopted. Specimens were loaded to failure at a constant cross-head rate of 1 [mm/min] using a 2 [kN] load cell.

## Result and discussion

Results from the misalignment analysis by FFNM on representative plates is given in Figure 2. Due to the intense processing pressure, both UD2/100% and UD14/100% plaques exhibited an in-plane fiber waviness in the range [-5° 5°] (Figure 2a), demonstrating that this material was able to flow even at 100% areal coverage. Apparent in-plane waviness was much less pronounced in TW6/100% and TW10/100% plaques, being consistently in the range of [0° 2°] as reported in Figure 2b. In-plane fiber waviness of the outer layer of UD2/100%/TW1/100%/SMC2/80% hybrid plaques was consistently in the range of [-3° 2°], demonstrating the effectiveness of TW fabric in preventing lateral motion of UD fibers during processing.
<figure><img src="image_505.png" /><figcaption>Figure 2: Fiber misalignment analysis of manufactured plates measured by FFNM technique. a) UD2/100% b) TW6/100% c) UD2/100%/TW1/100%/SMC2/80%</figcaption></figure>

Results from mechanical tests are summarized in Figure 3. Good repeatability was found for the elastic and tensile properties of discontinuous laminates with hybridization. The coefficient of variations (CV) was below 10% for the elastic modulus and below 5% for the ultimate tensile strength of UD and TW samples. However, the ultimate tensile strength of SMC specimens had a high scatter, with a CV of nearly 20%. The benefits of hybridization in stabilizing both elastic and ultimate tensile strength were confirmed by the low CV value of 6% for UD2/100%/TW1/100%/SMC2/80% samples. Comparable results were obtained in flexural tests. CV values for both UD and TW were below 5% for both flexural modulus and ultimate stress. On the other hand, the SMC samples demonstrated significant variability with a CV of 38% in modulus. The repeatability of results for hybrid UD2/100%/TW1/100%/SMC2/80% samples depended on the loading conditions of the continuous reinforcement.

<figure><img src="image_506.png" /><figcaption>Figure 3: Mechanical test results. a) tensile strength; b) tensile modulus; c) flexural strength d) flexural modulus</figcaption></figure>
<figure><img src="image_516.png" /><figcaption>Figure 2: Fiber misalignment analysis of manufactured plates measured by FFNM technique. a) UD2/100% b) TW6/100% c) UD2/100%/TW1/100%/SMC2/80%</figcaption></figure>

Results from mechanical tests are summarized in Figure 3. Good repeatability was found for the elastic and tensile properties of discontinuous laminates with hybridization. The coefficient of variations (CV) was below 10% for the elastic modulus and below 5% for the ultimate tensile strength of UD and TW samples. However, the ultimate tensile strength of SMC specimens had a high scatter, with a CV of nearly 20%. The benefits of hybridization in stabilizing both elastic and ultimate tensile strength were confirmed by the low CV value of 6% for UD2/100%/TW1/100%/SMC2/80% samples. Comparable results were obtained in flexural tests. CV values for both UD and TW were below 5% for both flexural modulus and ultimate stress. On the other hand, the SMC samples demonstrated significant variability with a CV of 38% in modulus. The repeatability of results for hybrid UD2/100%/TW1/100%/SMC2/80% samples depended on the loading conditions of the continuous reinforcement.

<figure><img src="image_517.png" /><figcaption>Figure 3: Mechanical test results. a) tensile strength; b) tensile modulus; c) flexural strength d) flexural modulus</figcaption></figure>
In the specimens where the long-fiber reinforcement was placed in tension during the test, the CVs of flexural modulus and tensile strength were found to be 7% and 5%, respectively. On the other hand, in specimens where the long-fiber reinforcement was placed in compression, the CVs of flexural modulus and tensile strength increased to 15% and 10%, respectively. Both tensile and flexural modulus of hybrid laminate was comparable to the ones of TW (Figure 3b, Figure 3c) although their resistance was slightly lower. A significative improvement in bending strength can be achieved by positioning the continuous reinforcement in traction, which contributes in increase the resistance by an additional 20% (Figure 3c). Representative stress-strain curves are shown in Figure 3. UD sets the upper bound for stiffness and strength among base laminates, while SMC sets the lower bound. The typical behavior of the stress-strain curve in tensile for all laminate is mostly linear except for the last 10% of the strain for UD2/100% material where the curves usually tend to be non-linear (Figure 4a) due to progressive failure of the material. In a previous study by the authors [10], it was observed that the initial macroscopic failure of UD2/100% specimens occurred in the areas where the maximum absolute value of fiber waviness was detected. Stress-strain curves for SMC3/80%, TW6/100%, and UD2/100%/TW1/100%/SMC2/80% samples in tensile tests exhibited consistent linear behavior until failure. Stress strain curves from flexural test highlight the different failure mechanisms of the materials used. Both UD14/100% and TW10/100% samples exhibited a linear region followed by a sudden load drop before abrupt failure. The strain field resulting from flexural tests of SMC8/80% samples was qualitatively spoken homogeneous in the beginning of loading. Afterwards a number of load-drops followed by load recovery was consistently observed across all samples, indicating the progressive damage of the material. It can be argued that these phenomena are strictly related to the heterogeneous microstructure of the SMC material, capable to provide a stress redistribution after onset of damage such as fiber debonding or cracking of the matrix [13].

<figure><img src="image_507.png" /><figcaption>Figure 4: Representative stress-strain curves from mechanical characterization. a) tensile b) flexural</figcaption></figure>

When short fibers are loaded in tension, these phenomena are observable also at intermediate loads on UD2/100%/TW1/100%/SMC2/80% samples as visible in Figure 4b. At higher loads, the hybrid laminate gradually fails, with several consecutive load drops preceding the final failure of the sample. When short fibers are loaded in compression, no load drop is visible on UD2/100%/TW1/100%/SMC2/80% samples. Up to the point of first failure, the stress-strain curve for this configuration overlaps with the TW samples. However, due to the inherent capacity of SMC to internally redistribute stresses, this configuration provides additional load-bearing capacity as visible by the largest deformation at break shown in Figure 4b. The DIC analysis provided a great insight in understanding the failure mechanisms of hybrid UD2/100%/TW1/100%/SMC2/80% samples during loading. On the SMC side of the specimen,
In the specimens where the long-fiber reinforcement was placed in tension during the test, the CVs of flexural modulus and tensile strength were found to be 7% and 5%, respectively. On the other hand, in specimens where the long-fiber reinforcement was placed in compression, the CVs of flexural modulus and tensile strength increased to 15% and 10%, respectively. Both tensile and flexural modulus of hybrid laminate was comparable to the ones of TW (Figure 3b, Figure 3c) although their resistance was slightly lower. A significative improvement in bending strength can be achieved by positioning the continuous reinforcement in traction, which contributes in increase the resistance by an additional 20% (Figure 3c). Representative stress-strain curves are shown in Figure 3. UD sets the upper bound for stiffness and strength among base laminates, while SMC sets the lower bound. The typical behavior of the stress-strain curve in tensile for all laminate is mostly linear except for the last 10% of the strain for UD2/100% material where the curves usually tend to be non-linear (Figure 4a) due to progressive failure of the material. In a previous study by the authors [10], it was observed that the initial macroscopic failure of UD2/100% specimens occurred in the areas where the maximum absolute value of fiber waviness was detected. Stress-strain curves for SMC3/80%, TW6/100%, and UD2/100%/TW1/100%/SMC2/80% samples in tensile tests exhibited consistent linear behavior until failure. Stress strain curves from flexural test highlight the different failure mechanisms of the materials used. Both UD14/100% and TW10/100% samples exhibited a linear region followed by a sudden load drop before abrupt failure. The strain field resulting from flexural tests of SMC8/80% samples was qualitatively spoken homogeneous in the beginning of loading. Afterwards a number of load-drops followed by load recovery was consistently observed across all samples, indicating the progressive damage of the material. It can be argued that these phenomena are strictly related to the heterogeneous microstructure of the SMC material, capable to provide a stress redistribution after onset of damage such as fiber debonding or cracking of the matrix [13].

<figure><img src="image_518.png" /><figcaption>Figure 4: Representative stress-strain curves from mechanical characterization. a) tensile b) flexural</figcaption></figure>

When short fibers are loaded in tension, these phenomena are observable also at intermediate loads on UD2/100%/TW1/100%/SMC2/80% samples as visible in Figure 4b. At higher loads, the hybrid laminate gradually fails, with several consecutive load drops preceding the final failure of the sample. When short fibers are loaded in compression, no load drop is visible on UD2/100%/TW1/100%/SMC2/80% samples. Up to the point of first failure, the stress-strain curve for this configuration overlaps with the TW samples. However, due to the inherent capacity of SMC to internally redistribute stresses, this configuration provides additional load-bearing capacity as visible by the largest deformation at break shown in Figure 4b. The DIC analysis provided a great insight in understanding the failure mechanisms of hybrid UD2/100%/TW1/100%/SMC2/80% samples during loading. On the SMC side of the specimen,