Limiting role of crystalline domain orientation on the modulus and strength of aramid fibers

Highlights

• Stress conditioning increases the crystallite orientation and elastic modulus of Kevlar but not the tensile strength.

• Regardless of initial crystallite orientation, all conditioned Kevlar fibers reached the same final modulus.

• Tensile strength is shown to be independent of length for Kevlar fibers as short as 200 μm.

• Very short Kevlar fibers (200 μm) showed failure initiation at the skin or the skin-core interface.

• The shear strength of the skin-core interface limits the fiber tensile strength.

Abstract

The evolution of crystalline domain orientation vs. mechanical properties of aramid fibers under mechanical loading was investigated for initial crystalline domain orientations between 16.7° and 9.7°. The latter resulted in a broad range of longitudinal moduli between 66 GPa and 119 GPa but tensile strengths in the narrow range of 3.5–4.0 GPa. Mechanical conditioning up to 90% of the fiber tensile strength increased the initial modulus converging to 100 GPa which corresponds to a stable crystalline domain orientation of 11.6°, while the unloading modulus at stresses near the tensile strength converged to ∼165 GPa which approaches the theoretical modulus of 220 GPa. On the contrary, the tensile strength remained unchanged with increasing crystalline domain orientation, and was shown to be independent of the fiber gauge length, thus implying that failure is not due to flaws obeying weakest link statistics. Instead, short gauge length tests (200 μm) showed failure initiation at the fiber skin followed by crack propagation at the skin-core interface, resulting in extrusion of the fiber core, which points to the skin-core interface as an important factor limiting the tensile strength of this class of fibers.

Introduction

Aramid (aromatic polyamide) fibers are characterized by high tensile strength and modulus [[1], [2], [3], [4]] due to the high molecular orientation along the fiber axis, which results in loading of the covalent bonds in the poly(p-phenylene terephthalamide) (PPTA) backbone [2,5,6]. The paraorientation of the aromatic rings in the PPTA macromolecule is due to the trans-conformation, while rotation of the PPTA chain around the phenyl-carbonyl and phenyl-nitrogen bonds is restricted due to the interactions between sequential phenyl and amine segments, thus resulting in a rigid rod-like microstructure [3,[7], [8], [9], [10]]. Paraorientation of the aromatic rings not only results in rigid rod-like chains but also enables efficient packing in a highly crystalline structure [2].

During fabrication, PPTA is dissolved in ≥99.8% sulphuric acid (H₂SO₄) to form the spinning solution which is isotropic at low concentrations but forms liquid crystalline domains as the concentration increases [11,12]. Fibers are spun at the solubility limit of PPTA at 80 °C with as many nematic liquid crystallites as possible for easier alignment of the molecular structure along the longitudinal axis of the fiber [11,13]. The fibers are obtained from the spinning solution in a dry-jet wet spinning process, during which the solution is pushed through a spinneret into an air gap, finally entering a coagulation bath which is kept at low temperature for optimal fiber formation [14,15]. The PPTA solution experiences shear flow through the spinneret resulting in alignment of the nematic crystallites. Moreover, the air gap allows for further alignment by stretching the filaments through extensional flow. It has been shown that the orientation imparted at the spinneret and the air gap is maintained during coagulation as the high rate of coagulation prevents molecular relaxation, while the loss of H₂SO₄ results in lateral shrinkage [13,16].

Understanding the microstructural effects on the mechanical behavior of oriented paracrystalline fibers has been critical in the historical evolution of high performance fibers. Aramid fibers are assumed to be comprised of orthotropic domains which are not perfectly oriented along the fiber axis but follow an orientation distribution. The axial elastic response has been shown to be predicted with sufficient accuracy via aggregate theory by accounting for the total deformation of individual domains and their orientation distribution along the fiber [[17], [18], [19]]. Further studies improved on this idea by assuming oblong domains that satisfy continuity of deformation between crystalline domains [20] and applied finite deformation theory to improve on the accuracy of predictions [21,22]. Other models have further built on the fibrillar and microvoid nature of aramid fibers by employing a supramolecular pleated sheet structure [23] to explain the coupling between extensional and shear stresses [24].

Although our understanding of the processes occurring in aramid fibers at small elastic stresses is fairly thorough, the deformation processes taking place in the non-linear regime of the stress-strain curves are less explored. At high longitudinal tension, the misorientation of crystalline domains (crystallites) induces significant local shear stresses which would result in domain rotation. However, it is not clear whether post-fabrication attempts to increase the domain orientation would also increase the tangent modulus, and to what extent such an increase would be retained upon unloading. Although limited experimental data from experimental PPTA fibers have shown an increasing tensile strength with orientation [25] this correlation is not apparent in Kevlar^® fibers which are the commercial form of aramid fibers. Rather, Kevlar^® fibers with quite different initial crystalline domain orientations and moduli have not demonstrated correspondingly different tensile strength values.

This study comes to address the question of evolution of molecular orientation with mechanical loading, starting with different initial conditions, and its effect on the fiber tensile strength. The correspondence between the tensile stiffness and strength and the induced orientation is explored by loading individual fibers to different strain levels. Towards this goal, we take advantage of high quality commercial aramid fibers, Kevlar^®, with different as-fabricated orientation distributions to deduce correlations between the initial orientation and its evolution with mechanical drawing, and to further explore the limits of controlling mechanical strength via crystalline domain orientation. Finally, we investigate the presence of strength limiting critical flaws as they may be manifested in fiber gauge length size effects on the tensile strength in pristine and oriented aramid fibers, reaching gauge lengths as short as 200 μm. We report on new and revealing findings from experiments conducted with the shortest fiber gauge lengths (200 μm) tested to date.