Spinline Cooling as a Determinant of Crystalline Structure and Mechanical Properties in Melt-Spun UHMWPE/HDPE Blend Fibers

Abstract

This study investigates the influence of cooling rates on the structural evolution and mechanical properties of ultra-high-molecular-weight polyethylene/high-density polyethylene fibers by systematically varying cooling media from ambient air (f1) to room-temperature water (f5). A significant structural inversion was observed between the as-spun and drawn fiber stages: while slow cooling (f1) promotes thermodynamic crystallization to form large, stable grains and maximum initial crystallinity (54%), rapid quenching (f5) effectively “freezes” the molecular chains in a low-crystallinity, highly orientable precursor state by suppressing thermal relaxation. During subsequent hot-drawing, the quenched samples (f5) exhibited a superior response to tensile stress, achieving the highest maximum draw ratio due to reduced crystalline obstacles and enhanced chain mobility. This enables efficient stress-induced crystallization, leading to near-perfect crystal orientation (f_c > 0.95) and a dense microfibrillar framework. Consequently, the fiber performance trends reversed, with f5 achieving peak tensile strength (1.33 GPa) and modulus, whereas f1 remained limited by its rigid thermal history. These findings highlight that rapid quenching is essential for developing high-performance fibers by preserving a precursor structure that maximizes the potential of stress-induced crystallization.

1. Introduction

Ultra-high-molecular-weight polyethylene (UHMWPE) fibers exhibit high strength and modulus, low density, and outstanding corrosion resistance, which underpin their widespread use in protective systems and aerospace applications [1,2]. Their mechanical performance ultimately derives from the extremely high molecular weight, which enables the development of highly oriented load-bearing chains along the fiber axis [3]. Yet this same molecular architecture also results in strong chain entanglement, giving rise to extremely high melt viscosity and pronounced viscoelasticity. As a consequence, conventional melt processing is prone to flow instabilities and melt fracture during draw-down and quenching/solidification, limiting the scalable manufacture of high-performance UHMWPE fibers [4,5]. A range of approaches have therefore been pursued to improve the processability of UHMWPE, including compositional modification and process design. Blending UHMWPE with high-density polyethylene (HDPE) is widely regarded as a practical route that balances industrial feasibility with mechanical performance [6,7,8]. The addition of HDPE can markedly improve melt flow by reducing viscosity and elastic response, thereby mitigating melt fracture and enabling stable melt spinning [9]. Wang et al. [10] reported that, with an optimized UHMWPE/HDPE ratio followed by thermal drawing, the tensile strength of the blended fibers can exceed 1.1 GPa; the improvement was attributed to HDPE-assisted disentanglement and the formation of a more regular and compact crystalline structure. In addition to polymer-polymer blending, the incorporation of nanofillers such as carbon nanofibers or silica has been introduced to further tune rheology and crystallization kinetics, providing additional—though often incremental—gains in mechanical performance [11,12]. However, although HDPE modification enables good melt spinnability of the UHMWPE system, the structure and mechanical properties of the resulting blended fibers still depend on further optimization of key processing conditions during spinning.

In the multi-step process of melt spinning high-performance polyethylene fibers, cooling and solidification of the extruded melt immediately downstream of the spinneret constitute a critical stage that governs the structural evolution of the nascent filament and its subsequent drawability [13]. Li et al. [14] reported that, for solvent-containing spinning filaments, solidification after the spinneret can proceed via air cooling, accompanied by rapid solvent evaporation. Subsequent studies have shown that, for materials with sufficiently high crystallization temperature and crystallization rate, post-spinneret air cooling can sustain stable melt spinning without inducing melt instability or surface defects, while yielding fibers with smooth surfaces and uniform structures [15]. Further studies [16] have shown that as melt-spun nascent filaments exit the spinneret, they undergo controlled laminar air cooling, which acts as a quenching step. This process enhances solidification and contributes to improved structural uniformity in the resulting fibers, and ref. [17] further demonstrates that air-cooling conditions not only govern filament cooling and drawing behavior, but also play a decisive role in the structural development of semicrystalline polymer fibers, where structural heterogeneity directly determines property variations. Liu et al. [18] employed room-temperature water quenching to prepare non-isothermal crystallization samples for investigating shear-induced disentanglement in UHMWPE/HDPE blends. In the preparation of UHMWPE fibers, Yan et al. [19] employed a cooling water bath at 5–15 °C placed downstream of the spinneret to obtain primarily cooled filaments. The water bath rapidly lowered the melt temperature and induced solidification. Seyhan et al. [20] showed that during melt spinning, introducing a liquid bath between the spinneret and the take-up rollers strengthens fiber–liquid interaction and raises the drag force on the filament. These effects together improve molecular chain orientation, leading to significantly enhanced mechanical properties in polyethylene fibers. The bath also promotes favorable interfacial effects, which lower interfacial energy and improve fiber wettability. To further clarify the distinct roles of different cooling media in regulating the structure and properties of UHMWPE-based materials, systematic comparisons among multiple cooling strategies have been undertaken. Kanagaraj et al. [21] demonstrated that the cooling method significantly influences the performance of UHMWPE. Air-cooled samples exhibited higher storage modulus and thermal stability, whereas water-cooled samples showed superior strength and toughness. Yu et al. [22] further noted that both the cooling medium and cooling rate significantly influence the crystalline morphology of polyethylene. Natural air cooling promotes the formation of larger crystalline grains, whereas rapid air cooling and water quenching refine grain size and accelerate crystallization. It was also observed that water cooling may introduce some additional dielectric loss.

Despite broad recognition of the importance of cooling conditions in regulating the structure and properties of polyethylene materials, systematic investigations focusing on melt-spun UHMWPE/HDPE blended fibers remain relatively limited. In particular, direct comparisons between different cooling media and conditions—such as water baths at varying temperatures and air cooling under natural versus forced convection—are scarce, and their effects on nascent filament structure evolution, drawability, and final mechanical performance are not yet well understood. Accordingly, the present study aims to systematically elucidate the effects of different cooling conditions on the structure and mechanical properties of melt-spun UHMWPE/HDPE blended fibers. Nascent fibers were prepared under five cooling conditions (ambient air, forced air, ambient water, 50 °C water bath, and 90 °C water bath), followed by thermal drawing. The evolution of thermal behavior, morphology, crystalline structure, orientation, and mechanical properties was characterized using DSC, SEM, WAXD/SAXS, sonic velocity, and tensile testing, with the objective of clarifying the trade-off between fiber strength and drawability and providing guidance for optimizing the overall performance of UHMWPE/HDPE blended fibers. Importantly, the results demonstrate a cooling-dependent inversion between the as-spun and drawn states, where rapid cooling yields a low-crystallinity yet highly deformable as-spun structure that facilitates stress-induced crystallization and ultimately leads to enhanced mechanical performance after drawing, highlighting the dominant role of processing-defined chain mobility and orientability rather than initial crystallinity alone.

2. Materials and Methods

2.1. Materials

The UHMWPE resin employed in this work was GUR4012 (Celanese Corporation, Nanjing, China), characterized by a weight-average molecular weight (Mw) of 1.5 × 106.

The HDPE material selected for blending was grade 5000S supplied by Yanshan Petrochemical Company (Beijing, China), which exhibited a melt flow index (MFI) of 0.9 g/10 min and a weight-average molecular weight of 1.64 × 105.

2.2. Preparation of UHMWPE/HDPE Blend Fibers

HDPE pellets and UHMWPE powder were first homogenized in a high-speed mixer at ambient temperature, operating at a rotational speed of 1440 rpm for 20 min. The blending ratio of UHMWPE to HDPE was fixed at 6:4 by mass. Subsequently, the blended materials were pelletized using a co-rotating twin-screw extruder with a screw diameter of 20 mm and a length-to-diameter ratio (L/D) of 45. During extrusion, the screw rotation speed was maintained at 200 r/min, while the temperatures of the five heating zones were set in the range of 120–300 °C.

The polymer melt was processed through a custom-designed single-screw extruder for extrusion. Prior to winding, the extrudate was subjected to different cooling environments, and the corresponding cooling conditions and sample designations are summarized in Table 1. The barrel temperature of the extruder was controlled within the range of 160–300 °C. A spinneret with a hole diameter of 1.2 mm was employed, and the spinneret temperature was maintained at 310 °C. Subsequently, the as-spun filaments were hot-drawn at 90 °C to produce fibers with tailored mechanical performance. A schematic diagram of the process is shown in Figure 1.

Table 1. Experimental Conditions and Expected Cooling Rates.

Sample ID	Cooling Medium	Temperature & Conditions	Cooling Intensity
f1	Ambient air	Natural Convection	Extremely Low
f2	Ambient air	Forced Convection (Air Cooling)	Low
f3	Water	90 °C	Medium
f4	Water	50 °C	High
f5	Water	~25 °C (Room Temperature)	Very High

Figure 1. Schematic diagram of the fiber preparation process.

2.3. Analytical Methods

The thermal behavior of both as-spun filaments and drawn fibers was characterized using differential scanning calorimetry (DSC2, Mettler Toledo, Greifensee, Switzerland). Each measurement was performed on samples with a mass of approximately 5 mg. DSC scans were conducted at a constant heating rate of 10 °C/min over a temperature interval ranging from 25 to 250 °C. All tests were carried out under a nitrogen atmosphere to prevent oxidative effects. The degree of crystallinity of the samples was calculated according to Equation (1):

$$X_c={\displaystyle \frac{\mathsf{\Delta }{H}_m}{\mathsf{\Delta }{H}_m^0}}\times 100\mathrm{\%}$$

(1)

where $X_c$ represents the degree of crystallinity determined from DSC analysis, $∆H_m$ denotes the melting enthalpy of the sample, and ${∆H}_m^0$ corresponds to the melting enthalpy of a fully crystalline material, which was taken as 293 J/g according to previously reported values in the literature [23,24].

Wide-angle X-ray diffraction (WAXD) measurements of the (110), (200), and (020) crystallographic planes were performed using a Bruker AXS D8 Advance diffractometer (Bruker AXS GmbH, Karlsruhe, Germany) equipped with Cu Kα radiation (λKαl = 1.542 Å). Data were collected over a 2θ range of 5–50° at a scanning rate of 10°/min. The diffraction behavior of the (002) plane was examined separately using an EMPYREAN X-ray diffractometer (Malvern Panalytical, Malvern, UK) with Cu Kα radiation, operating over a 2θ range of 70–78° at a step size of 0.013°. For azimuthal analysis, the 2θ angle was fixed at the maximum intensity position (21.4°), and azimuthal scans were conducted from −90° to 90° at a scanning speed of 0.1°/min. The two-theta position was calibrated using the diffraction peak of standard silicon powder. The crystallite size (L_hkl) perpendicular to the corresponding lattice planes was calculated using the Scherrer equation:

$$L_{hkl}={\displaystyle \frac{K\lambda }{Hcos\theta }}$$

(2)

where λ denotes the X-ray wavelength, θ represents the Bragg angle, and H refers to the full width at half maximum of the diffraction peak, expressed in radians. The orientation degree of the crystalline phases in the fibers was subsequently evaluated using Equations (3) and (4):

$$f_c={\displaystyle \frac{1}{2}}(3{cos}^2\phi -1)$$

(3)

$${cos}^2\phi ={\displaystyle \frac{{\int }_{-\frac{\pi }{2}}^{\frac{\pi }{2}}{I}_{hkl}sin\phi {cos}^2\phi d\phi }{{\int }_{-\frac{\pi }{2}}^{\frac{\pi }{2}}{I}_{hkl}sin\phi d\phi }}$$

(4)

where f_c denotes the orientation degree of the crystalline phase, taking values of 0, 1, and −1/2 for randomly oriented crystals, perfectly aligned crystals, and crystals exhibiting perfect normal orientation relative to the drawing direction, respectively. The I_hkl distribution was obtained through azimuthal scanning, where I_hkl represents the diffraction intensity of a given crystallographic plane and φ is the corresponding azimuthal angle.

The sound velocity orientation of the samples was measured using an SSY-III sound velocity orientation instrument (Beijing Hydro-Tech Marine Technology Co., Ltd., Beijing, China). The sound velocity was determined by the double-length measurement method. The corresponding sound velocity orientation factor of the samples was calculated according to Equation (5):

$$f_s=(1-{\displaystyle \frac{C^2}{C_m^2}})\times 100$$

(5)

where $C$ represents the sound velocity of the sample, and $C_m$ denotes the sound velocity of polyethylene fibers with random molecular orientation, which was taken as 1.65 km/s according to the literature [25].

The mechanical performance of the fibers was assessed using a YG001A-1 electronic fiber tensile tester (Hongda Fangyuan Electric Co., Ltd., Suzhou, China). Fiber specimens with a gauge length of 20 mm were tested at a constant crosshead speed of 10 mm/min. The fiber diameters were determined using a SMART biological microscope (Drawell, Chongqing, China). For each sample, at least five specimens were tested, and the results are presented as mean values accompanied by the corresponding standard deviations.

The morphological features of the fibers were characterized via scanning electron microscopy (S4700 SEM, Hitachi Ltd., Tokyo, Japan) at an accelerating voltage of 20 kV.

3. Results and Discussion

3.1. SEM Studies

The surface morphologies of the as-spun and drawn fibers under varying cooling conditions were characterized by SEM, as illustrated in Figure 2. In the as-spun stage (Figure 2a,c,e,g,i), a subtle increase in surface roughness is observed as the cooling rate intensifies from f1 (ambient air cooling) to f5 (water quenching). Remarkably, after subsequent hot-drawing (Figure 2b,d,f,h,j), the initial morphological discrepancies vanish, with all samples—particularly the initially rougher f5—transforming into uniform, smooth, and streamlined microfibrillar structures. This evolution from a rough precursor to a smooth fiber provides evidence of structural reorganization during drawing.

Figure 2. SEM images of samples: (a) f1 as-spun filament; (b) f1 fiber; (c) f2 as-spun filament; (d) f2 fiber; (e) f3 as-spun filament; (f) f3 fiber; (g) f4 as-spun filament; (h) f4 fiber; (i) f5 as-spun filament; (j) f5 fiber.

3.2. Sound Velocity Orientation Studies

As shown in Figure 3, the molecular chain orientation degree (f_s) for both as-spun filaments and fibers increases progressively from f1 to f5. This trend is intrinsically linked to the increasing cooling gradient, which governs the kinetic competition between stress-induced orientation during processing and the spontaneous thermal relaxation of polymer chains. For f5, the exceptionally high cooling rate provided by the room-temperature water bath induces an immediate quenching effect. This process rapidly vitrifies the polymer matrix, effectively “freezing” the molecular chains in their extended conformation and preventing the entropic retraction (disorientation) that typically occurs in the melt state. In contrast, the lower f_s values observed in f1 result from the prolonged residence time in a semi-molten state during ambient air cooling. The sustained high-temperature environment provides sufficient chain mobility for the relaxation rate to outweigh the freezing rate, leading to the dissipation of the orientation acquired during spinning.

Figure 3. Molecular chain orientation degree (f_s) of the samples: (a) as-spun filaments; (b) fibers.

3.3. DSC Analyses

To further investigate the thermal behavior and crystalline evolution regulated by cooling kinetics, DSC analysis was performed on both as-spun filaments and drawn fibers. As illustrated in Figure 4a, the as-spun filaments exhibit a melting peak temperature (T_m) shift and varying enthalpy areas, where the thermodynamic-driven crystallization order follows f1 > f2 > f3 > f4 > f5. Specifically, sample f1, subjected to the slowest cooling rate in ambient air, displays the highest melting enthalpy and a sharper endothermic peak, indicating the formation of more stable, thicker lamellae and higher initial crystallinity. In sharp contrast, sample f5 shows a broader and shallower melting peak at lower temperature, caused by the quenching effect. This result indicates that rapid heat transfer restricts chain folding and yields a crystalline state with smaller, less perfect grains. A notable structural inversion, however, occurs in the fiber after hot-drawing (Figure 4b). The melting peaks of the drawn fibers become noticeably sharper and shift to higher temperatures. Sample f5 now shows the most prominent endothermic peak and the highest T_m. This transformation confirms that the highly oriented yet low-crystallinity precursor structure in f5 effectively promotes stress-induced crystallization during drawing. Consequently, f5 surpasses f1 in both crystalline perfection and thermal stability.

Figure 4. DSC curves of the samples: (a) as-spun filaments; (b) fibers.

As shown in Figure 5, the crystallinity evolution from as-spun filaments to drawn fibers reveals a pronounced trend reversal. For as-spun filaments (Figure 5a), the crystallinity decreases monotonically with increasing cooling rates (f1 > f2 > f3 > f4 > f5), a behavior primarily governed by the competition between thermodynamic driving forces and kinetic limitations. In f1 (ambient air cooling), the low thermal conductivity of air allows the polymer to remain within its optimal crystallization temperature range for an extended duration, granting molecular chains ample time to organize into stable crystalline domains. In contrast, f5 (water quenching) subjects the melt to severe supercooling, rapidly vitrifying the disordered molecular network and yielding the lowest initial crystallinity. A striking transition occurs at the fiber stage (Figure 5b), where the crystallinity order is completely reversed (f5 > f4 > f3 > f2 > f1). This ‘structural inversion’ is intrinsically linked to the distinct precursor states: although f5 exhibits the lowest initial crystallinity (X_c,DSC = 40.23%), its rapid quenching preserves a highly mobile, orientable amorphous network. During hot-drawing, these mobile chains undergo efficient stress-induced crystallization (SIC), facilitating the development of a highly aligned microfibrillar structure. Conversely, the high initial crystallinity and stable grain structures of f1 act as physical constraints that resist molecular reorganization. According to the Peterlin model [26], such pre-existing crystalline obstacles restrict plastic deformation and prevent the formation of the highly ordered framework observed in quenched samples.

Figure 5. Crystallinity values of the samples: (a) as-spun filaments; (b) fibers.

3.4. TG and DTG Analyses

The thermal stability of the fibers was evaluated via TGA and DTG, as shown in Figure 6. All samples exhibited a single-step degradation at about 470 °C, corresponding to the random scission of polyethylene chains. For the as-spun filaments, thermal stability followed the order of f1 > f2 > f3 > f4 > f5, with the air-cooled f1 displaying the highest maximum decomposition temperature (T_max). This aligns with the DSC results, where slow cooling (f1) promoted a more robust thermodynamic crystalline network (54.01% crystallinity) that effectively restricted thermal chain motion. However, a significant “structural inversion” occurred after hot-drawing, with the stability reversing to f5 > f1. Despite its lower initial stability, the water-quenched f5 precursor-characterized by high chain mobility in DSC-facilitated more efficient stress-induced crystallization during drawing. This resulted in a more compact and ordered molecular packing in the final f5 fiber, leading to a higher T_max and narrower DTG peak compared to f1, where coarse initial crystals hindered uniform structural reorganization. Consequently, the rapid quenching process (f5) ultimately yields fibers with superior thermal resistance.

Figure 6. TG and DTG curves of samples: (a) TG curves of as-spun filaments; (b) DTG curves of as-spun filaments; (c) TG curves of fibers; (d) DTG curves of fibers.

3.5. WAXD Analyses

Figure 7 displays the WAXD curves of as-spun filaments and fibers, with the inset tables listing the corresponding FWHM values. For the as-spun filaments (Figure 7a), a clear trend is observed where the (110) peak width increases from f1 (0.26°) to f5 (0.32°). This peak broadening indicates a reduction in the lateral crystallite size, as the rapid quenching in f5 limits the time available for molecular chains to organize into larger, more perfect crystalline domains. In the fiber stage (Figure 7b), all samples exhibit significantly higher diffraction intensities compared to the as-spun state. However, the f5 fiber continues to show a broader (110) peak (FWHM of 0.57°) than the f1 fiber (FWHM of 0.43°). This suggests that the initial disordered state of the f5 precursor leads to a final crystalline structure characterized by more refined crystalline units.

Figure 7. WAXD curves of the samples: (a) as-spun filaments; (b) fibers.

The degree of crystallinity derived from WAXD (X_c,WAXD) was compared with the DSC results (X_c,DSC) in Table 2. For all samples, X_c,WAXD values are consistently lower than X_c,DSC by approximately 4–10%, a common discrepancy in semi-crystalline polymers. Despite this difference in sensitivity, both techniques independently verify the “structural inversion” phenomenon: slow cooling (f1) maximizes thermodynamic crystallinity in the as-spun stage, whereas rapid quenching (f5) facilitates superior stress-induced crystallization during drawing, ultimately yielding the highest final crystallinity.

Table 2. Comparison of Degrees of Crystallinity, Determined by DSC (X_c,DSC) and WAXD (X_c,WAXD), for As-spun Filaments and Fully Drawn Fibers.

Sample	As-Spun Filaments Xc,DSC (%)	As-Spun Filaments Xc,WAXD (%)	Fully Drawn Fibers Xc,DSC (%)	Fully Drawn Fibers Xc,WAXD (%)
f1	54.01	45.9	58.32	52.5
f2	52.55	44.7	58.65	53.1
f3	45.36	38.6	60.34	55.5
f4	41.74	35.5	63.79	59.3
f5	40.23	34.2	65.52	61.6

The evolution in microstructure, specifically in crystallinity and grain size, indicates a fundamental shift in crystallization mechanisms between the as-spun and fiber stages. As observed in Figure 5a and Figure 8a, the crystallinity and grain size of the as-spun filaments consistently decrease in the order of f1 > f2 > f3 > f4 > f5. This trend is consistent with the kinetics of thermal crystallization. Specifically, the slow air cooling for f1 results in a narrow undercooling degree, which inhibits nucleation but promotes crystal growth. Consequently, molecular chains have sufficient time to reorganize into thick lamellae and form large, well-developed grains. In contrast, the rapid quenching of f5 in the 25 °C water bath creates a high degree of supercooling, which drives burst-like heterogeneous nucleation. The subsequent immediate temperature drops below the crystallization threshold effectively “freezes” these nuclei, preventing significant growth. Consequently, f5 exhibits the lowest initial crystallinity and the smallest grain size among all samples.

Figure 8. Grain size analysis of the samples: (a) as-spun filaments; (b) fibers.

The subsequent drawing process induces a clear structural inversion, as most evident in the fiber morphology shown in Figure 5b and Figure 8b. The fiber crystallinity and axial grain size (notably L₀₀₂) reverse to follow the order f5 > f4 > f3 > f2 > f1. This inversion results from stress-induced crystallization. The high cooling rate of f5 preserves a disordered but highly orientable amorphous network despite its low initial crystallinity. This enables the molecular chains to be forcefully extended and incorporated into the crystal lattice under the maximum draw ratio, thereby generating a high volume fraction of new crystals. The intense tensile stress simultaneously refines the grains radially and elongates them along the fiber axis. This is manifested in f5 as the longest coherent scattering length in the axial direction, with the structure evolving into fine, needle-like micro-fibrillar crystals. Conversely, the pre-existing large grains in f1 resist such reorganization, leading to inferior final orientation and limited axial elongation under the same processing conditions.

The thermal history and subsequent mechanical deformation are decisive for the structural evolution, as jointly evidenced by the pronounced variations in crystal orientation (f_c), crystallinity, and grain size. As shown in Figure 9, the crystal orientation follows the order f5 > f4 > f3 > f2 > f1 for both as-spun filaments and fibers. At the as-spun stage, the high cooling rate of f5 effectively preserves the flow-induced precursor orientation, which acts as a template for the subsequent “oriented crystallization.” In contrast, the slow cooling of f1 permits sufficient molecular relaxation, resulting in a more isotropic crystalline arrangement characterized by spherulite formation and significantly lower f_c values. During the drawing process, the pre-aligned amorphous chains in f5 are efficiently integrated into the crystal lattice along the fiber axis, enabling it to achieve near-perfect orientation (f_c > 0.95). In contrast, the presence of unoriented crystalline regions in f1 resists realignment under stress, thereby resulting in a lower final orientation degree.

Figure 9. Crystal orientation (f_c) of samples: (a) as-spun filament; (b) fibers.

This orientation behavior is inherently linked to the inversion of crystallinity and grain size trends observed between the as-spun and fiber stages. Initially, as-spun samples exhibit a crystallinity and grain size hierarchy of f1 > f2 > f3 > f4 > f5 (Figure 5a and Figure 8a). Although the slow cooling of f1 yields low undercooling, favoring the formation of large, thick lamellae and high thermal crystallinity, a remarkable transition occurs during high-ratio drawing: f5 ultimately surpasses f1 in both final crystallinity and axial grain size (L₀₀₂), as shown in Figure 5b and Figure 8b. This inversion is driven by stress-induced crystallization, in which the highly oriented amorphous network in quenched f5 is efficiently converted into fine, needle-like micro-fibrillar crystals under intense tensile stress. In contrast, the larger, pre-existing grains in f1 cannot achieve comparable axial elongation, retaining a shorter and broader morphology that ultimately limits the final fiber performance. A schematic illustration of this process is presented in Figure 10.

Figure 10. Schematic illustration of the crystalline grain evolution within the fibers.

3.6. Mechanical Properties

The mechanical properties of the drawn fibers (Figure 11) exhibit a clear dependence on the spinline cooling rate, with the tensile strength and initial modulus following the sequence of f5 > f4 > f3 > f2 > f1. The f5 fiber, originating from the most rapidly quenched precursor, achieves the highest tensile strength of 1.33 GPa. This superior performance is further elucidated by the correlation between crystallinity (X_c) and tensile strength, as shown in Figure 12. A strong positive relationship is observed, confirming that the crystalline domains act as the primary load-bearing elements within the oriented microfibrillar structure. Notably, although f1 possessed the highest initial crystallinity in its as-spun state, its final mechanical performance was restricted by the coarse crystalline obstacles formed during slow cooling. In contrast, the low-crystallinity ‘optimized precursor’ of f5 facilitated more efficient stress-induced crystallization during the hot-drawing process. This led to the highest final X_c, which, combined with the near-perfect crystal orientation (f_c), significantly enhanced the efficiency of stress transfer and resulted in the most robust mechanical integrity among all samples.

Figure 11. Mechanical properties and drawability of the fiber samples: (a) maximum drawing ratio; (b) tensile strength (MPa); (c) initial modulus (GPa); and (d) elongation at break (%).

Figure 12. Correlation between the degree of crystallinity (X_c) and the tensile strength of the fully drawn UHMWPE/HDPE blend fibers (f1–f5).

To further elucidate the underlying mechanism, a comprehensive structure–property relationship is summarized in Figure 13. In the f1 regime, slow cooling permits extensive chain relaxation and promotes the formation of stable, bulky lamellae, which function as “geometric constraints”, thereby limiting effective load transfer and subsequent molecular orientation. In contrast, the rapid quenching in f5 preserves a “highly orientable precursor” characterized by a disordered but pre-aligned amorphous network. During hot-drawing, this unique precursor allows for the maximum conversion of amorphous chains into highly extended-chain crystals (shish-kebab or microfibrillar structures), as evidenced by the significant increment in L₀₀₂ and f_c. Ultimately, this transition from a kinetically trapped disordered state to a stress-induced highly ordered state plays a critical role in achieving superior tensile strength and modulus in UHMWPE/HDPE blend fibers.

Figure 13. Schematic diagram of the structure-property relationship governing the mechanical performance of the fibers.

4. Conclusions

This study demonstrates that the cooling rate of as-spun filaments is the primary factor governing the final mechanical performance of high-performance fibers through a “structure–precursor” mechanism. Overall, an inverse relationship between initial crystalline structure and drawability was established. Slow cooling (f1) favors the formation of higher initial crystallinity and larger grains, which ultimately hinder effective drawing. In contrast, rapid water quenching (f5) preserves a low-crystallinity, highly orientable amorphous network, enabling efficient stress-induced crystallization during subsequent deformation. Specifically, slow-cooled filaments exhibit higher crystallinity and coarser crystalline grains in the as-spun state, whereas rapidly quenched filaments show suppressed crystal growth but significantly improved drawability. Upon drawing, the rapidly cooled samples display more efficient stress-induced crystallization, higher crystalline orientation, and refined microfibrillar structures. By maximizing the draw ratio, f5 achieves a near-perfect crystal orientation (f_c > 0.95) and a dense microfibrillar structure, resulting in a tensile strength of 1.3 GPa. Importantly, the final mechanical performance is not dictated solely by the absolute crystallinity, but by the synergistic effects of crystallinity, crystalline orientation, and crystal morphology that are established during stress-induced crystallization. These findings provide a theoretical basis for optimizing industrial spinning processes by balancing the kinetic competition between freezing and relaxation rates to enhance fiber performance.