Effect of aging treatment on the precipitation behavior and mechanical properties of Mg-Gd-Zr alloy

ABSTRACT

This study conducts a comprehensive investigation into the effect of isothermal aging on the microstructural evolution and mechanical properties of a Mg-8.2Gd-0.5Zr (wt.%) alloy. The alloy was subjected to solution treatment followed by aging at three distinct temperatures: 200 °C, 225 °C, and 250 °C for durations up to 300 hours. The age-hardening response was characterized by Vickers microhardness measurements, revealing a strong dependence on temperature, with the highest peak hardness of 135 HV achieved at 200 °C after 96 hours. Detailed microstructural analysis using X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM) established the precipitation sequence as Supersaturated Solid Solution (SSSS) → β″ (D019) → β′ (bco) → β (FCC). The peak-aged condition was associated with a high number density of fine, lenticular βʹ precipitates on the prismatic planes of the α-Mg matrix. Tensile testing demonstrated that the peak-aged alloy exhibited a remarkable increase in strength, with a yield strength of 290 MPa and an ultimate tensile strength of 380 MPa, albeit with a significant reduction in ductility. Fracture surface analysis revealed a transition from ductile dimple fracture in the solution-treated state to brittle cleavage fracture in the peak-aged condition. A quantitative analysis of strengthening mechanisms confirms that precipitation hardening via the Orowan mechanism is the dominant contributor to the alloy’s strength, accounting for approximately 75% of the increase in yield strength (≈70–80% when input uncertainties are considered). These findings provide a systematic understanding of the structure-property relationships in Mg-Gd-Zr alloys, offering a basis for optimizing heat treatments to achieve desired mechanical performance.

Keywords:
Orowan strengthening; Precipitate morphology; Fracture mechanism; Isothermal heat treatment; Microstructure-property correlation

1. INTRODUCTION

The persistent demand for lightweight, high-performance structural materials in the aerospace, automotive, and defense sectors has intensified research into advanced metallic alloys [¹]. Magnesium (Mg) alloys, being the lightest of all structural metals, are exceptionally attractive due to their high specific strength and stiffness, offering significant potential for weight reduction and improved energy efficiency [²]. However, the widespread application of conventional Mg alloys has been historically constrained by their inherent limitations, including relatively low absolute strength, poor formability at room temperature, and inadequate mechanical performance at elevated temperatures, which stem from the hexagonal close-packed (HCP) crystal structure with its limited number of active slip systems.

Alloying has emerged as the most effective strategy to overcome these deficiencies. Among various alloying elements, rare-earth (RE) elements, particularly Gadolinium (Gd), have garnered considerable attention [³]. Gd exhibits substantial solid solubility in the Mg matrix at elevated temperatures, which decreases sharply upon cooling, creating a strong thermodynamic driving force for precipitation hardening [⁴]. This age-hardening response is a powerful mechanism for enhancing the strength of Mg alloys [⁵]. Concurrently, the addition of Zirconium (Zr) is a cornerstone of modern high-strength Mg alloy design. Zr is recognized as the most potent grain refiner for aluminum-free Mg alloys, effectively reducing the grain size through heterogeneous nucleation. This grain refinement not only improves strength and hardness via the Hall-Petch mechanism but also enhances ductility and processability, making the material more amenable to subsequent forming operations [¹]. The combination of a potent precipitation strengthener like Gd with an effective grain refiner like Zr represents a sophisticated metallurgical design strategy, allowing for the decoupling and independent optimization of baseline properties and age-hardened performance [⁶].

The strengthening effect in Mg-Gd based alloys is primarily derived from the controlled precipitation of metastable phases during aging. The generally accepted precipitation sequence proceeds from a supersaturated solid solution (SSSS) through several intermediate phases: SSSS → β″ → βʹ → β. The metastable β″ (D019 structure) and βʹ (base-centered orthorhombic, bco) phases are coherent or semi-coherent with the α-Mg matrix and act as highly effective obstacles to dislocation motion, leading to significant increases in strength [⁴]. The β′ phase, in particular, is widely regarded as the principal strengthening precipitate in these systems [⁷]. While this general sequence is known, the precise kinetics, morphology, crystallography, and spatial distribution of these precipitates are acutely sensitive to the aging parameters, most notably temperature and time.

A systematic and comparative investigation is required to fully elucidate how aging temperature governs the microstructural evolution and, in turn, dictates the full spectrum of mechanical properties. This includes not only hardness and strength but also ductility and the operative fracture mechanisms, which are critical for structural reliability. Therefore, the objective of the present study is to conduct a comprehensive investigation into the effect of isothermal aging at three distinct temperatures—200 °C, 225 °C, and 250 °C—on the precipitation behavior and mechanical properties of a model Mg-8.2Gd-0.5Zr (wt.%) alloy. This work aims to establish clear, quantitative relationships between processing parameters, microstructural features, and mechanical performance, culminating in a quantitative assessment of the contributions from various strengthening mechanisms.

2. MATERIALS AND METHODS

2.1. Alloy synthesis and composition

The experimental alloy with a nominal composition of Mg-8.2Gd-0.5Zr (in weight percent) was prepared for this investigation. High-purity raw materials, including Mg (99.98%), Gd (99.95%), and a Mg-30Zr (wt.%) master alloy, were used. The materials were melted in a graphite crucible within an electrical resistance induction furnace. To prevent oxidation and ignition of the molten magnesium, a protective atmosphere of carbon dioxide (CO₂) and sulfur hexafluoride (SF₆) in a volumetric ratio of 100:1 was maintained throughout the melting process [⁸]. The melt was held at a temperature of 780 °C for 30 minutes to ensure the complete dissolution and homogenization of the alloying elements. Subsequently, the melt was cooled to a pouring temperature of 730 °C and cast into a cylindrical steel mold preheated to 250 °C. The final chemical composition of the as-cast ingot was determined using Inductively Coupled Plasma-Atomic Emission Spectrometry (ICP-AES). The measured composition, presented in Table 1, was found to be in close agreement with the nominal target.

Table 1
Chemical composition of the investigated Mg-8.2Gd-0.5Zr alloy (wt.%).

ELEMENT	Gd	Zr	Mg
Composition (wt.%)	8.15 ± 0.05	0.48 ± 0.02	Balance

2.2. Heat treatment

Specimens machined from the as-cast ingot underwent a multi-stage heat treatment process. First, a homogenization treatment was performed at 520 °C for 16 hours in an air-circulating furnace to eliminate dendritic segregation and ensure a chemically uniform matrix [⁹]. This was immediately followed by a solid solution treatment at 525 °C for 8 hours, designed to dissolve the Gd-containing secondary phases into the α-Mg matrix, thereby forming a supersaturated solid solution (SSSS). After the solution treatment, the specimens were rapidly quenched into warm water maintained at approximately 80 °C to retain the supersaturated state and minimize thermal stresses [⁴]. This solution-treated and quenched condition is hereafter referred to as T4.

Following the T4 treatment, specimens were subjected to isothermal aging. The aging was conducted in a thermostatically controlled molten salt bath furnace to ensure precise temperature control and rapid heat transfer [¹⁰]. Three aging temperatures were selected for this study: 200 °C, 225 °C, and 250 °C. For each temperature, specimens were aged for various durations, ranging from 0.5 to 300 hours, to systematically capture the entire aging sequence from the under-aged to the severely over-aged state [¹¹].

2.3. Microstructural characterization

The microstructure of the alloy at various stages of heat treatment was characterized using a suite of analytical techniques. For optical microscopy (OM) and scanning electron microscopy (SEM), specimens were prepared using standard metallographic procedures. This involved grinding with silicon carbide (SiC) papers up to 4000 grit, followed by polishing with 3 µm and 1 µm diamond suspensions to achieve a mirror-like surface finish. The polished specimens were then etched with a solution consisting of 5 g picric acid, 5 mL acetic acid, 10 mL deionized water, and 80 mL ethanol to reveal the grain structure and secondary phases [⁴]. SEM analysis was performed on a FEI Quanta 250 FEG-SEM.

Phase identification was carried out using X-ray diffraction (XRD) on a Bruker D8 Advance diffractometer with Cu-Kα radiation (λ = 0.15406 nm). The XRD scans were conducted over a 2θ range of 20° to 90° with a step size of 0.02°.

Quantitative analysis of grain size, grain boundary characteristics, and crystallographic texture was performed using Electron Backscatter Diffraction (EBSD). The EBSD analysis was conducted on the same SEM, equipped with an Oxford Instruments Nordlys detector. Specimens for EBSD were subjected to a final polishing step using a colloidal silica suspension to ensure a high-quality surface finish.

Detailed investigation of the nanoscale precipitates was performed using transmission electron microscopy (TEM). Thin foils for TEM were prepared by mechanically grinding 3 mm discs to a thickness of approximately 50 µm. These foils were then thinned to electron transparency using a Gatan PIPS II ion milling system with a low-angle, low-energy final polishing step to minimize surface damage. The TEM analysis was conducted on a JEOL JEM-2100F field emission TEM operating at an accelerating voltage of 200 kV. This instrument was equipped with a High-Angle Annular Dark-Field (HAADF) detector, enabling Z-contrast imaging for the unambiguous identification of the Gd-rich precipitates. Precipitate morphology, size, distribution, and crystal structure were determined through a combination of bright-field (BF) imaging, selected area electron diffraction (SAED), and atomic-resolution HAADF-STEM imaging.

2.4. Mechanical property evaluation

The mechanical response of the alloy was evaluated through microhardness and uniaxial tensile testing. Vickers microhardness (HV) measurements were performed on the polished metallographic samples using a Mitutoyo HM-200 tester. A load of 200 gf (1.96 N) and a dwell time of 15 s were applied for each indentation. To ensure statistical accuracy, at least ten indentations were made for each specimen condition, and the average value was reported. These measurements were used to construct the age-hardening curves for each aging temperature [⁴].

Room-temperature uniaxial tensile tests were conducted to determine the 0.2% proof stress (yield strength, YS), ultimate tensile strength (UTS), and elongation to failure (EL). The tests were performed on an Instron 5985 universal testing machine in accordance with the ASTM E8/E8M standard. Flat, dog-bone shaped tensile specimens with a gauge length of 25 mm and a cross-section of 6 mm × 2 mm were machined from the heat-treated samples. The tests were carried out at a constant strain rate of 1 × 10⁻³ s⁻¹ [¹²]. For each condition, at least three specimens were tested to ensure the reliability and reproducibility of the results.

Finally, the fracture surfaces of the tested tensile specimens were examined using the SEM to identify the dominant fracture modes and mechanisms under different aging conditions. This fractographic analysis provides crucial insight into the relationship between microstructure and ductility.

3. RESULTS AND DISCUSSION

3.1. Initial microstructure of the solution-treated alloy

The microstructure of the Mg-8.2Gd-0.5Zr alloy after solution treatment at 525 °C for 8 hours and subsequent quenching (T4 condition) serves as the baseline for evaluating the effects of aging. Figure 1 presents the characterization of this initial state. The optical micrograph in Figure 1a reveals a homogeneous and fully recrystallized microstructure. The grains are equiaxed, indicating that the solution treatment was effective in eliminating the as-cast dendritic structure.

Figure 1
Microstructure of the Mg-8.2Gd-0.5Zr alloy in the solution-treated (T4) condition. (a) Optical micrograph showing the general equiaxed grain structure. (b) EBSD Inverse Pole Figure (IPF) map confirming the equiaxed morphology and random texture. (c) EBSD-derived grain size distribution histogram, indicating an average grain size of 45 µm.

Further quantitative analysis by EBSD, shown in the Inverse Pole Figure (IPF) map in Figure 1b, confirms the equiaxed grain morphology. The random color distribution within the IPF map signifies a weak and random crystallographic texture, which is typical for a cast and solution-treated alloy that has not undergone significant plastic deformation. The corresponding grain size distribution histogram, presented in Figure 1c, shows a relatively uniform distribution with a statistically calculated average grain size of approximately 45 µm. This fine grain structure is a direct result of the addition of Zr, which acts as a powerful grain refiner in Mg alloys by providing a high density of potent heterogeneous nucleation sites during solidification [¹³]. This refined grain structure is expected to provide a beneficial baseline contribution to both strength and ductility through the Hall-Petch mechanism.

3.2. Age-hardening behavior

The age-hardening response of the T4-treated alloy was investigated by isothermal aging at 200 °C, 225 °C, and 250 °C. The evolution of Vickers microhardness as a function of aging time is presented in Figure 2. The initial hardness of the SSSS alloy was approximately 75 HV. All three aging temperatures produced a characteristic age-hardening profile, consisting of an initial increase in hardness to a peak value, followed by a gradual decrease during over-aging. The aging kinetics and the peak hardness achieved were strongly dependent on the aging temperature [¹⁴]. At 200 °C, the alloy exhibited the most potent hardening response, although with the slowest kinetics. The hardness increased steadily, reaching a maximum peak value of approximately 135 HV after a prolonged aging time of 96 hours. Increasing the aging temperature to 225 °C significantly accelerated the process. At this temperature, a peak hardness of approximately 125 HV was attained after only 24 hours. At the highest aging temperature of 250 °C, the kinetics were further accelerated, with a peak hardness of 110 HV reached in just 8 hours. However, this was followed by a more rapid softening during the over-aging stage.

Figure 2
Age-hardening curves of the Mg-8.2Gd-0.5Zr alloy, plotting Vickers hardness (HV) as a function of aging time at 200 °C, 225 °C, and 250 °C.

This behavior highlights a critical inverse relationship between aging temperature and achievable peak hardness in this alloy system. The attainment of maximum hardness is governed by the formation of a dense dispersion of fine, coherent or semi-coherent precipitates that effectively impede dislocation motion [¹⁵]. The precipitation process involves a competition between nucleation and growth, both of which are thermally activated. At the lower temperature of 200 °C, the high supersaturation of Gd in the Mg matrix provides a strong thermodynamic driving force for nucleation. This condition favors the formation of a very large number of small precipitate nuclei. Concurrently, the low temperature limits atomic diffusion, slowing the growth and coarsening of these precipitates. The result is a microstructure containing a high number density of finely dispersed precipitates, leading to the highest peak hardness [¹⁶]. Conversely, at the higher temperature of 250 °C, the driving force for nucleation is reduced while atomic diffusion is significantly enhanced. This leads to the formation of fewer, larger nuclei that coarsen rapidly, resulting in a less dense and coarser precipitate distribution that is less effective for strengthening, and thus a lower peak hardness [¹⁷].

3.3. Phase identification and microstructural evolution

To understand the microstructural origins of the observed hardening behavior, detailed phase analysis and microscopic investigations were conducted. Figure 3 shows the XRD patterns of the alloy in the T4 condition, the peak-aged condition (T6, aged at 200 °C for 96 hours), and a severely over-aged condition (aged at 200 °C for 300 hours). The pattern for the T4 sample is dominated by sharp diffraction peaks corresponding to the α-Mg matrix, confirming the formation of a nearly single-phase SSSS. In the T6 peak-aged condition, the pattern still primarily shows α-Mg peaks [¹⁸]. The principal strengthening precipitates (β″ and β′) are metastable and coherent or semi-coherent with the matrix, and their volume fraction and crystal structure do not produce distinct, high-intensity diffraction peaks that are easily distinguishable from the matrix background. However, in the severely over-aged sample, new peaks emerge that can be indexed to the equilibrium β phase (Mg₅Gd), which has a face-centered cubic (FCC) crystal structure [¹⁹]. The appearance of this phase coincides with the softening observed in the hardness curves, confirming that over-aging involves the transformation of metastable strengthening phases into a stable, non-strengthening phase [²⁰]. A weak, broad feature at 2θ ≈ 22–23° is ascribed to amorphous scattering from the glass holder and does not correspond to any Mg–Gd or Mg–Zr crystalline reflection; the phase analysis and conclusions are unchanged.

Figure 3
XRD patterns for the Mg-8.2Gd-0.5Zr alloy in the solution-treated (T4), peak-aged (T6, 200 °C for 96 h), and over-aged (200 °C for 300 h) conditions.

The evolution of the precipitate microstructure during aging at 200 °C was investigated in detail using TEM and HAADF-STEM, as this condition yielded the highest strength. The precipitation sequence was determined to be SSSS → β″ → β′ → β, consistent with previous studies on Mg-Gd systems [¹¹]. In the early, under-aged stage (16 hours at 200 °C), as shown in Figure 4, fine, dot-like precipitates begin to form. The BF-TEM image (Figure 4a) shows a high density of these nascent precipitates. The corresponding SAED pattern along the α zone axis (Figure 4c) exhibits faint streaks along the <10–10>α directions, which is characteristic of the formation of the precursor β″ phase. The β″ phase has a D019 crystal structure and forms as small, coherent particles [²¹]. The HAADF-STEM image (Figure 4b) provides Z-contrast, where the Gd-rich β″ precipitates appear as bright dots against the darker Mg matrix.

Figure 4
TEM characterization of the alloy in the under-aged condition (200 °C for 16 h). (a) Bright-field (BF) TEM image showing fine, dot-like precipitates. (b) HAADF-STEM image revealing the distribution of Gd-rich β'' precipitates. (c) SAED pattern along the α zone axis with characteristic streaks.

Upon reaching the peak-aged condition (96 hours at 200 °C), the microstructure is dominated by a very high number density of the metastable βʹ phase, as shown in Figure 5. These precipitates have a lenticular or plate-like morphology and are uniformly distributed throughout the matrix. The BF-TEM image (Figure 5a) illustrates their dense and fine dispersion. When viewed along the [¹¹–²⁰]α zone axis (Figure 5b), the βʹ plates are observed to lie with their broad faces on the {10–10}α prismatic planes of the Mg matrix. At peak hardness, these precipitates have an average length of approximately 50 nm and a thickness of 10–15 nm. The SAED pattern taken along the α zone axis (Figure 5c) shows distinct superlattice reflections at the 1/2 positions between the matrix spots, confirming the base-centered orthorhombic (bco) crystal structure of the βʹ phase [²², ²³]. A high-resolution HAADF-STEM image of a single βʹ precipitate (Figure 5d) reveals its semi-coherent interface with the matrix. This dense array of βʹ precipitates provides the maximum impediment to dislocation motion, corresponding to the peak hardness observed in Figure 2.

Figure 5
TEM characterization of the alloy in the peak-aged condition (200 °C for 96 h). (a) BF-TEM image showing a high number density of fine, lenticular β' precipitates. (b) HAADF-STEM image along the [¹¹–²⁰]α zone axis showing β' plates on prismatic planes. (c) SAED pattern along the α zone axis with superlattice reflections confirming the bco structure of β'. (d) High-resolution HAADF-STEM image of a single β' precipitate.

In the over-aged condition (300 hours at 200 °C), the microstructure undergoes significant coarsening, as shown in Figure 6. The βʹ precipitates have grown larger and their number density has decreased (Figure 6a). More significantly, the equilibrium β phase begins to form, typically as coarser, rod-like particles, often nucleating at grain boundaries. The formation of these incoherent, stable particles is accompanied by the development of a Precipitate-Free Zone (PFZ) adjacent to the grain boundaries (Figure 6b). A PFZ is a region depleted of solute (Gd) due to diffusion to the grain boundary to form the coarse β precipitates [²⁴]. This region is soft and acts as a site for preferential strain localization, contributing significantly to the loss of strength and hardness during over-aging [²⁵].

Figure 6
TEM characterization of the alloy in the over-aged condition (200 °C for 300 h). (a) BF-TEM image showing coarsened β' precipitates and larger, rod-like equilibrium β phase particles. (b) Image showing the formation of a Precipitate-Free Zone (PFZ) along a grain boundary.

To understand the lower peak hardness at higher aging temperatures, the precipitate microstructures at the respective peak-aged conditions were compared. Figure 7 shows the microstructure after aging at 225 °C for 24 hours, while Figure 8 shows the microstructure after aging at 250 °C for 8 hours. In both cases, the primary strengthening phase is still the βʹ precipitate. However, a direct comparison with the 200 °C peak-aged condition reveals a clear trend. As the aging temperature increases, the number density of the βʹ precipitates visibly decreases, while their average size (both length and thickness) increases [²⁶]. The precipitates in the sample aged at 225 °C (Figure 7) are noticeably larger and more widely spaced than those aged at 200 °C. This trend is even more pronounced in the sample aged at 250 °C (Figure 8), where the precipitate distribution is significantly coarser and less dense. This microstructural coarsening directly explains the reduction in peak hardness with increasing aging temperature. A coarser, less dense dispersion of precipitates provides fewer and weaker obstacles to dislocation motion, resulting in a lower strengthening effect [²⁷]. This observation reinforces the conclusion that achieving maximum strength in this alloy requires aging conditions that favor high nucleation rates and slow growth, which are characteristic of lower aging temperatures.

Figure 7
TEM characterization of the alloy at peak hardness for aging at 225 °C (24 h). (a) BF-TEM image showing the distribution of β' precipitates. (b) HAADF-STEM image showing the size and spacing of the β' plates for comparison.

Figure 8
TEM characterization of the alloy at peak hardness for aging at 250 °C (8 h). (a) BF-TEM image showing a coarser distribution of β' precipitates. (b) HAADF-STEM image showing significantly larger and more widely spaced β' plates.

3.4. Mechanical properties and structure-property correlation

The room-temperature tensile properties of the alloy in the T4, T6 (peak-aged at 200 °C), and over-aged (200 °C, 300 h) conditions were measured to correlate the observed microstructures with mechanical performance. The results are summarized in Table 2. The T4-treated alloy exhibited a modest yield strength of 130 MPa, an ultimate tensile strength of 240 MPa, and excellent ductility with an elongation to failure of 15.2%. This behavior is characteristic of a soft, solid-solution strengthened, and fine-grained material [²⁸].

Table 2
Room temperature mechanical properties of the Mg-8.2Gd-0.5Zr alloy under various aging conditions.

CONDITION	YS (MPa)	UTS (MPa)	ELONGATION (%)
T4 (Solution-Treated)	130 ± 5	240 ± 8	15.2 ± 0.8
T6 (Peak-Aged at 200 °C)	290 ± 6	380 ± 7	3.1 ± 0.4
Over-Aged (200 °C, 300 h)	215 ± 7	310 ± 9	8.5 ± 0.6

After peak aging at 200 °C for 96 hours, the alloy demonstrated a dramatic improvement in strength. The yield strength increased by over 120% to 290 MPa, and the ultimate tensile strength reached 380 MPa. This substantial strengthening is a direct consequence of the dense and uniform dispersion of βʹ precipitates throughout the matrix. These semi-coherent precipitates act as powerful obstacles to dislocation glide, requiring a much higher stress to initiate and sustain plastic deformation. This strengthening can be primarily described by the Orowan bowing mechanism, where dislocations are forced to bow out between the hard precipitates [²⁹]. However, this remarkable increase in strength came at the cost of a severe loss in ductility. The elongation to failure for the T6 sample plummeted to just 3.1%. This pronounced strength-ductility trade-off is a common feature in high-strength, precipitation-hardened alloys. The underlying mechanism for this embrittlement is intense strain localization. In the T6 state, the dense βʹ precipitates effectively pin dislocations on the primary basal slip planes. This leads to the formation of large dislocation pile-ups at the precipitate-matrix interfaces, creating high local stress concentrations. Since non-basal slip and cross-slip are difficult to activate in Mg at room temperature, these stresses cannot be easily relieved [³⁰]. Eventually, the stress concentration becomes high enough to initiate microcracks, either by decohesion at the interface or by cleavage fracture of the matrix along the basal plane where dislocations have piled up. Once initiated, these microcracks can propagate rapidly through the brittle, precipitate-filled matrix, leading to premature failure and low macroscopic ductility.

In the over-aged condition, the mechanical properties were intermediate. The yield strength and UTS decreased to 215 MPa and 310 MPa, respectively, while the elongation recovered to 8.5%. This behavior is consistent with the coarsened microstructure. The larger, more widely spaced precipitates are less effective at impeding dislocation motion, leading to lower strength [³¹]. Concurrently, the presence of soft PFZs and the reduced overall obstacle density allow for more homogeneous deformation, resulting in a partial recovery of ductility. To complement the tensile property data listed in Table 2, illustrative stress–strain curves were constructed based on the measured yield strength, ultimate tensile strength, and elongation values (Figure 9). These curves highlight the characteristic trade-off between strength and ductility, particularly the dramatic loss of plasticity in the T6 peak-aged condition compared to the T4 solution-treated alloy. In the over-aged state, the partial recovery of elongation is also evident.

Figure 9
Stress–strain curves of the Mg-8.2Gd-0.5Zr alloy at room temperature under three aging conditions (T4 solution- treated, T6 peak-aged, and over-aged), ductility after over-aging.

3.5. Fracture behavior

The fracture surfaces of the tensile specimens were examined by SEM to identify the failure mechanisms, providing further evidence for the observed ductility trends. The fractographs are presented in Figure 10.

Figure 10
SEM fractographs of the tensile specimens. (a) Ductile fracture surface with fine dimples in the T4 condition. (b) Brittle, transgranular cleavage fracture in the T6 peak-aged condition. (c) Mixed-mode fracture with both cleavage facets and dimples in the over-aged condition.

The T4-treated sample (Figure 10a) exhibits a fracture surface characterized by a dense population of fine, equiaxed dimples. This is the hallmark of a ductile fracture mechanism known as microvoid coalescence, which is consistent with the high elongation measured for this condition.

In stark contrast, the fracture surface of the T6 peak-aged sample (Figure 10b) is dominated by large, flat facets with distinct river patterns. This morphology is characteristic of transgranular cleavage fracture, a brittle failure mode. This observation directly supports the hypothesis of strain localization and microcrack-induced failure discussed previously [³²]. The high stresses generated by dislocation pile-ups at the βʹ precipitates lead to cleavage along crystallographic planes, resulting in the observed brittle behavior. The over-aged sample (Figure 10c) displays a mixed-mode fracture surface. It contains regions of cleavage facets, but also a significant presence of dimpled rupture. This mixed morphology reflects the intermediate ductility of the over-aged condition, where some regions fail in a brittle manner while others undergo ductile tearing.

3.6. Quantitative analysis of strengthening mechanisms

To provide a deeper understanding of the exceptional strength of the T6 peak-aged alloy, the contributions of the various operative strengthening mechanisms to the yield strength were quantitatively evaluated. The overall yield strength (σ_y) can be modeled as a linear superposition of several components [³³]:

$${\sigma }_{\text{y}}={\sigma }_{\text{0}}+{\sigma }_{\text{ss}}+{\sigma }_{\text{gb}}+{\sigma }_{\text{p}}$$

where σ₀ is the intrinsic lattice friction stress of pure Mg, σ_ss is the contribution from solid solution strengthening, σ_gb is the grain boundary strengthening contribution (Hall-Petch effect), and σ_p is the precipitation strengthening contribution.

Intrinsic Lattice Resistance (σ₀): This is the baseline resistance of the crystal lattice to dislocation motion. For pure magnesium, a value of σ₀ = 11 MPa is commonly used [³⁴].

Solid Solution Strengthening (σ_ss): After precipitation, some Gd atoms remain dissolved in the α-Mg matrix. This contribution is estimated using the relation Δσ_ss = k_ss^⋅c^1/2, where k_ss is a constant for Gd in Mg and c is the atomic concentration of residual Gd in the matrix. Based on literature values and an estimated residual Gd concentration of ~1.5 wt.%, the contribution is calculated to be approximately 35 MPa [³³].

Grain Boundary Strengthening (σ_gb): This contribution is calculated using the Hall-Petch equation: Δσ_gb = kHP^⋅d^−1/2, where kHP is the Hall-Petch coefficient and d is the average grain size. Using the measured grain size of d = 45 µm and a literature value of kHP = 164 MPa·µm^1/2 for Mg alloys, the grain boundary strengthening is calculated to be approximately 24 MPa [³⁴].

Precipitation Strengthening (σ_p): This is expected to be the dominant strengthening mechanism. It is estimated using the Orowan-Ashby equation for non-shearable precipitates:

$$\Delta {\sigma }_{\text{p}}=\text{π}\cdot 1-\text{v}\text{M}\cdot \text{0}\text{.4}\cdot \text{G}\cdot \text{b}\text{L}-\text{2rln}\text{(2r/b)}$$

where M is the Taylor factor (~2.5 for random texture), G is the shear modulus of Mg (17 GPa), b is the Burgers vector (0.32 nm), ν is Poisson’s ratio (0.35), L is the mean inter-precipitate spacing, and r is the mean radius of the precipitates. Using the dimensions of the βʹ precipitates measured from the TEM images in the T6 condition, the strengthening contribution from precipitation is calculated to be approximately 220 MPa.

The calculated contributions are summarized in Table 3 and visualized in Figure 11. The sum of the calculated strengthening contributions is 11 + 35 + 24 + 220 = 290 MPa. This calculated total yield strength is in excellent agreement with the experimentally measured value of 290 ± 6 MPa. This analysis quantitatively confirms that precipitation hardening is, by a significant margin, the most potent strengthening mechanism in the peak-aged Mg-8.2Gd-0.5Zr alloy, contributing approximately 75% of the total strength. A sensitivity analysis (see new ‘Parameter set and uncertainty bounds’ subsection) indicates an Orowan contribution in the range ≈70–80% when literature-supported variability in kHP, residual Gd in solution, elastic constants, and TEM-derived precipitate geometry/spacing are propagated.

Table 3
Calculated contributions of different strengthening mechanisms to the yield strength of the peak-aged (200°C, 96h) alloy.

STRENGTHENING MECHANISM	SYMBOL	CALCULATED CONTRIBUTION (MPa)
Intrinsic Lattice Resistance	σ0	11
Solid Solution Strengthening	σss	35
Grain Boundary Strengthening	σgb	24
Precipitation Strengthening	σp	220
Total Calculated Yield Strength	σy,calc	290
Experimental Yield Strength	σy,exp	290 ± 6

Figure 11
Bar chart illustrating the relative contributions of the different strengthening mechanisms to the total yield strength of the Mg-8.2Gd-0.5Zr alloy in the T6 peak-aged condition.

We report the parameter values and ranges used to compute each strengthening contribution and the associated sensitivity. For precipitation (Orowan) strengthening we treat β′ plates as non-shearable obstacles, consistent with established models for Mg–RE alloys; the model uses M = 2.5, G = 16.6–17.3 GPa, ν = 0.32–0.35, b = 0.321 nm, precipitate radius r inferred from TEM-measured thickness 10–15 nm (r≈5–7.5 nm), and inter-plate spacing L measured from TEM fields (intercept method; statistical details provided). For grain-boundary strengthening we adopt the Hall–Petch relation with d = 45 µm (EBSD) and kHP = 140–200 MPa·µm1/2 to reflect the reported variability with texture and testing conditions in Mg alloys. For solid-solution strengthening we use Δσss = ksscⁿ with n≈1/2–2/3 and matrix Gd content 1.2–1.8 wt.% after ageing, consistent with reported trends for Mg–Gd. Within these bounds the Orowan increment varies modestly (≈205–230 MPa), while σgb and σss vary more strongly (≈19–30 MPa and ≈25–45 MPa, respectively), yielding an Orowan fraction of ≈70–80% and a best estimate ≈75–76%. References documenting the adopted models and constants include Nie (review of precipitation hardening and β′ in Mg–RE), Gao–Chen–Han (solid-solution strengthening potency of Gd and combined mechanism accounting in Mg–Gd–Y–Zr), and Yu et al. (variability of kHP in Mg).

4. CONCLUSION

This investigation systematically characterized the effect of isothermal aging treatment on the microstructure and mechanical properties of a Mg-8.2Gd-0.5Zr alloy. The following conclusions can be drawn:

1. The Mg-8.2Gd-0.5Zr alloy demonstrates a potent age-hardening response. The kinetics of aging and the magnitude of the peak hardness are highly dependent on the aging temperature. Lower aging temperatures lead to slower kinetics but result in higher peak strength.

2. Aging at 200 °C for 96 hours produced the optimal combination of hardness and strength, achieving a peak hardness of 135 HV, a yield strength of 290 MPa, and an ultimate tensile strength of 380 MPa. This represents a more than two-fold increase in yield strength compared to the solution-treated condition.

3. The precipitation sequence during aging was determined to be SSSS → β″ (D019) → βʹ (bco) → β (FCC). The exceptional strength achieved in the peak-aged condition is unequivocally attributed to the formation of a high number density of fine, semi-coherent, plate-like βʹ precipitates, which are uniformly dispersed on the prismatic planes of the α-Mg matrix. Higher aging temperatures (225 °C and 250 °C) resulted in a coarser, less dense precipitate microstructure, leading to lower peak strength.

4. A classic strength-ductility trade-off was observed. The significant increase in strength upon peak aging was accompanied by a severe reduction in ductility, with elongation dropping from 15.2% to 3.1%. Fractographic analysis confirmed a corresponding transition in fracture mechanism from ductile microvoid coalescence in the T4 state to brittle transgranular cleavage in the T6 state.

5. A quantitative analysis of strengthening mechanisms successfully modeled the yield strength of the peak-aged alloy. The results confirmed that precipitation strengthening, as described by the Orowan mechanism, is the dominant strengthening contributor, accounting for approximately 76% of the total yield strength. This study provides a detailed framework for understanding and tailoring the mechanical properties of Mg-Gd-Zr alloys through the precise control of aging parameters.

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Publication Dates

• Publication in this collection
  20 Oct 2025
• Date of issue
  2025

History

• Received
  01 Aug 2025
• Accepted
  12 Sept 2025