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Article

Development of Zn-Reinforced Mg Matrix Composites via High Energy Ball Milling Duration: Impact on Mechanical Properties and Biodegradability

by
S. Bilal Çetinkal
1,*,
Emin Salur
1,
Gökhan Arıcı
1,
Ahmed Degnah
2,3,
Sayan Sarkar
4,5 and
Halit Sübütay
1
1
Department of Metallurgical and Materials Engineering, Selçuk University, Konya 42075, Turkey
2
Advanced Materials Technology Institute, King Abdulaziz City for Science and Technology, P.O. Box 6086, Riyadh 11442, Saudi Arabia
3
King Salman Center for Disability Research, Riyadh 11614, Saudi Arabia
4
Department of Materials Science, University of Utah, Salt Lake City, UT 84112, USA
5
Intel Corporation, Ronler Acres Campus, Hillsboro, OR 97124, USA
*
Author to whom correspondence should be addressed.
Coatings 2025, 15(5), 561; https://doi.org/10.3390/coatings15050561
Submission received: 27 March 2025 / Revised: 28 April 2025 / Accepted: 6 May 2025 / Published: 8 May 2025
(This article belongs to the Special Issue Advanced High Entropy Alloy Materials and Films: Properties and Applications)
Browse Figures
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Abstract

:
In this study, Zn-reinforced Mg matrix composite materials were produced via powder metallurgy by exposing them to ball milling at varying mechanical milling times. Following ball milling, the powders were cold-pressed under 600 MPa to obtain green compacts. The sintering process was carried out in a tube furnace under an argon atmosphere at 500 °C for 120 min. The effects of different milling times (2 h, 4 h, and 8 h) on particle and grain size, as well as the influence of sintering temperature and time on the microstructure, were investigated through SEM analysis. Phase evolution and changes in crystal planes occurring after ball milling were revealed by XRD analysis. SEM images show that Zn particles were homogeneously distributed within the matrix after 8 h of milling. Furthermore, it can be clearly stated that the highest hardness values were obtained from the samples produced after 8 h of milling. The sample group with the highest density, least mass loss, and lowest degradation rate was obtained from materials produced from 4 h ball milled powders. The intermetallic phase formed in the powder structure after 8 h of milling tends to reduce density and corrosion properties. The findings reveal that the addition of these alloys to pure Mg clearly enhances its hardness and density, while also imparting superior corrosion resistance. These combined improvements suggest that the developed materials hold strong potential for application in biomedical and clinical environments, where both mechanical strength and corrosion resistance are critical.

1. Introduction

Magnesium is preferred in many industries due to its specific properties. Its low density (1.738 g/cm3) and high specific strength make it valuable in the defense and transportation sectors; while its high thermal conductivity, excellent dimensional stability, good electromagnetic shielding, high damping capacity, good machinability, and ease of recycling make it useful in the automotive, computer, and aerospace industries. Additionally, due to its low specific weight and compatibility with metabolism, it is utilized in the healthcare sector [1,2,3,4,5,6].
Despite these advantageous properties, the limited strength of pure Mg restricts its independent use. For this reason, Mg alloys are formed with elements such as Al, Zn, Mn, Zr, Ca, Ag, Li, Ge, and As. In this way, strength, toughness, and corrosion resistance increase. In particular, Zn attracts attention as an additive that increases strength and corrosion resistance in studies [7]. In addition, it has been observed that corrosion resistance increases by adding Ca to the Mg–Zn alloy [8]; hardness and strength increase by adding Sr [9]; mechanical properties improve by adding hydroxy apatite (HA) [10]; corrosion resistance increases by adding Sn [11].
The Mg–Zn alloy is used in advanced technology fields to obtain lightweight, durable, and controlled dissolution materials. In biomedical applications such as bone screws, plates, and stents, the controlled dissolution rate of the alloy in biological environments is a critical factor. The ability of these alloys to gradually degrade within the body eliminates the need for secondary surgical interventions. In addition, Mg and Zn are minerals found in the human body and play an important role in metabolic processes. Research has identified at least 29 different elements, including both metals and non-metals, present in human tissues. Frieden (1985), who divided them into micro, trace, and ultra-trace elements according to the amount found in the tissues, stated that Mg and Zn are among the micro trace elements. Magnesium, whose daily requirement for humans is over 100 mg, is in Group II, and zinc, whose daily requirement is under 100 mg, is in Group III [12]. Given their biocompatibility and solubility, Mg and Zn have been extensively studied as biomaterials. A study on the biodegradability of Mg–Zn composites found that their mechanical properties were well-suited for implant applications. Additionally, intermetallic compounds formed during the casting process contributed to enhanced mechanical performance [13].
In parts with a complex shape, high-pressure casting is used instead of gravity casting. Thanks to high-pressure casting, low grain sizes are obtained in addition to intermetallics [14,15]. Another method that can be used in the production of complex-shaped parts is the powder metallurgy method. Powder metallurgy is a production method that allows the production of metallic materials with homogeneous and fine-grained microstructures [16,17]. Different mixing methods are used to ensure homogeneity in powder metallurgy. The most efficient of these methods is the mechanical alloying (MA) method. MA is a method of processing powder materials in a high-energy environment in a chamber. In this high-energy environment, the powders continuously enter the welding and fracture cycle. In this process, different material powders are mixed at the atomic scale and homogeneous alloys are obtained. MA is a method that provides homogeneous distribution of reinforcement nanoparticles in the matrix alloy material [18]. Unlike conventional mixing methods, MA not only ensures a homogeneous dispersion of reinforcement particles but also modifies the size and structure of the powders. In terms of the milling time effect on the nature of mechanical alloying, Toozandehjani et al. [19] found that prolonging milling time leads to detriment in the powder size as well as the porosity ratio. Ahamed and Senthilkumar [20] presented that increasing milling time conveys an evenly distributed nanoparticle matrix materials system which eventually enhances the hardness of the milled materials. MA is the origin for the formation of high energy generated in powders as a result of collision between grinding medium and powder particles which causes plastic deformation of powders, repeated fracture, and cold welding of particles [21,22,23].
The high mechanical properties of ball-milled products vary depending on the grinding parameters (powder ball ratio, grinding time, grinding speed, and grinding shape) [24]. Mechanical alloying has an important place in advanced technology and material science research because it can produce materials with nanostructured and intermetallic phases beyond the limits reached by traditional melting methods. In the literature, it is revealed that different sizes (micron and nano) and various types (oxide and carbide) of reinforcement materials can be distributed homogeneously in different matrix materials with the MA method and thus various nano-composite materials can be produced [13,25,26,27,28,29].
In general, a limited number of studies and information about Zn-reinforced Mg composites produced by MA technique is available in the known literature. Therefore, a demand arose to investigate mechanically alloyed Mg–Zn alloys with desired properties. In this study, the homogeneous distribution and biodegradability of 6% Zn reinforcement in Mg with different milling times and the production of the Zn-reinforced Mg composite material with the best biodegradability were investigated. The aim is to understand the effect of milling time on the microstructure, mechanical properties, and biodegradability after sintering.

2. Materials and Method

Production and Characterization of Powders

To fabricate Mg-based composites, pure magnesium powders (approximately 50 µm in size) were utilized as the matrix material, while zinc powders (99.9% purity and approximately < 5 µm in size) were employed as reinforcement elements. The magnesium matrix was reinforced by incorporating zinc powders at weight percentages of 6%.
Subsequent to the procurement of magnesium and zinc powders, preliminary trials were conducted using magnesium powder to identify the optimal Process Control Additive (PCA). Two of the most widely utilized PCA systems in the literature, stearic acid and methanol, were subjected to ball milling at various milling durations with a rotation speed of 300 rpm [30]. In accordance with the objectives of the study, where it was essential for the zinc powders to infiltrate the magnesium surface and form intermetallic compounds, the PCA system that facilitated the formation of a flake structure in the magnesium powder was preferred. This preference was based on the theoretical assumption that a more flake magnesium powder would increase the probability of zinc powders embedding into the surface. Scanning Electron Microscope (SEM) images revealed that stearic acid was the PCA system that induced a more flake structure in the magnesium powder across different milling times, while methanol facilitated cold welding and milling of the powders.
Upon identification of the appropriate PCA system, ball milling processes were conducted for the magnesium-based powders with wt% 6Zn reinforcement at various milling times. Ball milling parameters were established based on a combination of literature and experimental results. The determined parameters are presented in Table 1.
The composite powders produced and characterized within the content of the study were exactly weighed on a precision balance and placed into a mold with a 13 mm diameter.
Afterward, they were pressed under 600 MPa for 1 min using a cold pressing device, forming a uniaxially pressed specimen.
To enhance the high strength and density of the Mg–Zn alloys with different ball milling times, the sintering process, which is the second stage of the powder metallurgy production method, was applied. The sintering process was conducted in an atmosphere-controlled tube furnace under an argon gas atmosphere at a heating rate of 5 °C/min, maintaining a temperature of 500 °C for 2 h. For the microstructural observation of the produced metal matrix composite samples, the specimens were subjected to metallographic sample preparation procedures. The morphology of the composites was characterized using a Scanning Electron Microscope (SEM) (ZEISS, EVO LS10, Germany). The Brinell hardness tests of the samples were conducted using a hardness tester (Bulut Makina, Digirock, Turkey), with a 62.5 kg load, a 10 s dwell time, and a 2.5 mm diameter steel ball. The theoretical densities of the samples were calculated using the mixture rule equation (Equation (1)), and the experimental densities were measured based on the Archimedes principle (Equation (2)) (ASTM B962–17 standard) [31] with a precision balance (Precisa, XB 220A, Dietikon, Switzerland).
ρ c = w m ρ m + w r ρ r
In this Equation, ρm and ρr are the density of the matrix (Mg) and reinforcement materials (Zn), and wm and wr are the weight fraction of matrix and reinforcement materials, respectively.
ρ = w a w a − w w × ρ w
where w a the weight of samples in air, w w is the apparent weight of samples in water, ρ w is the density of the water, respectively.
Relative densities of materials were computed by mensural the material’s experimental density and dividing it by the material’s theoretical density.
Corrosion tests were conducted in accordance with the ASTM-G31-72 standard [32]. The behavior of Mg–6Zn alloys in Hank’s solution was determined by considering the weight percentage losses. Prior to the tests, the weights of each specimen were measured and listed individually. For each alloy, the precalculated amount of Hank’s solution was added to 50 mL beakers, and alloys with different milling time were placed into separate beakers. Every 24 h, the specimens were removed from the corrosion solution, rinsed out with deionized water to remove corrosion residues, and then thoroughly cleaned with a high-purity alcohol solution. After the cleaning and drying process, the weights of the specimens were measured individually. The solutions in each beaker were replaced, and the specimens were returned to the corrosion environment. Measurements were continued using this method for a period of 15 days. After the corrosion process, the weight percentage losses of each specimen were calculated using Equation (3).
T o t a l   W e i g h t   L o s s   P e r c e n t a g e % = I n i t i a l   w e i g h t   o f   s a m p l e − f i n a l   w e i g h t   o f   t h e   s a m p l e I n i t i a l   w e i g h t   o f   s a m p l e × 100
The annual degradation rates (mm/year) for all surfaces of the produced samples were calculated with the assistant of Equation (4) [33]. The terms expressed in the equation are defined as constant coefficient K = 8.76 × 104 (mm/year); W, weight loss of the sample (g); A, total surface area of the sample (cm2); T, corrosion exposure time (h); and D, sample density (gr/cm3) [34].
C o r r o s i o n   r a t e m m y e a r = K × W / A × D × T

3. Results and Discussion

3.1. Powder Morphology and Distribution of the Reinforcement Element in Different Ball Milling Durations

The images showing the initial particle size of the Mg and Zn elements are presented in Figure 1a,b. It can be observed from the relevant images that Mg powders have an irregular-shaped morphology, while Zn powders have a spherical shape. After 2 h of ball milling, the powders tend to flake due to the high energy resulting from PCA and the ball–powder–wall interaction (Figure 1c).
After 4 h of milling, the powder morphology changes to a form where small fragmented satellite particles, flakes, and platelet structures are seen together, due to the nature of the ball milling process (Figure 1d). As shown in the relevant image, after 8 h of milling, the particle shape changes and the particle size decreases. The mechanism responsible for the formation of the particle morphology observed in Figure 1e is severe plastic deformation. The excessive plastic deformation and work hardening resulting from the strong repeated collisions between the ball, powder, and wall during ball milling cause the flaky and platelet powder structure to break into smaller particles [35].
In the Mg–6% Zn material system, SEM mapping analysis was conducted to examine the distribution trend of Zn particles within the structure as a function of increasing ball milling time. Upon reviewing the SEM mapping images provided in Figure 2, it can be observed that, with increasing milling time, the Zn particles exhibit an irregular distribution in the powders obtained after 2 h of milling. In addition, based on the corresponding visuals, it can be concluded that Zn particles, initially smaller than 5 µm, tend to agglomerate after the milling process.
As a result of increasing the milling time to 4 h, it is evident from the corresponding images that the effectiveness of Zn agglomerates within the structure begins to decrease, and the distribution of the Zn element within the structure improves. Zn clusters, which were formed during the early stages of milling (2 and 4 h) and contributed to preventing homogeneous distribution, can be effectively broken down with the increase in milling time (8 h). It is observed that these broken-down Zn particles (regions shown in green in the relevant EDS mapping) exhibit a more regular distribution throughout the Mg matrix (which is represented in red). Particles that undergo increased ball–powder–wall interactions due to extended milling times tend to break and distribute more homogeneously within the structure [36]. It can be concluded that no random Zn cluster islands are formed within the structure following the 8 h milling process. Additionally, it can be stated that the 8 h milling period triggered the Zn element to adopt a more homogeneous distribution within the structure.

3.2. XRD Results Depending on the Various Milling Time in Mg–6ZN Powder System

Figure 3 shows the XRD spectra of pure Mg and the powders of the Mg–6Zn system subjected to ball milling for different times. The data in the graph reveal the hexagonal close-packed (HCP) structure of Mg (JCPDS card No. #35–0821) and Zn (JCPDS card No. #87-0713) elements. It is observed that the peaks corresponding to the Mg and Zn phases remain unchanged after 2 and 4 h of ball milling.
In addition, it is evident from the results that no new phases appear in the structure during these milling times. This clearly shown that Mg and Zn maintain their elemental stability during these milling durations, and the conditions required to trigger a chemical reaction are not met.
Although no new phases are observed during these two milling times (2 and 4 h), there are subtle changes in the peak profile parameters (peak position and intensity), which will be further discussed in subsequent sections. As the milling time increases, a new peak corresponding to the Mg7Zn3 intermetallic phase (JCPDS card No. #01-1185) emerges in the XRD spectrum [37].
Additionally, as the milling time gradually increases, the peak reflections associated with the Zn phase begin to decrease and ultimately disappear after 8 h of ball milling. Moreover, the peak intensity of the Mg7Zn3 intermetallic phase develops and becomes more prominent after 8 h of milling. There are two plausible explanations for the decrease in the peak intensity of the Zn phase with increasing milling time.
The first is that as the milling time increases from 2 to 4 h, Mg powders become flaky, allowing Zn to embed into the matrix or potentially dissolve. The second explanation is that the Zn phase acts as a precursor for the chemical reaction that forms the Mg7Zn3 intermetallic phase.
Therefore, after 8 h of milling, the Zn phase disappears, and the intermetallic phase becomes dominant. Based on these results, it can be concluded that a minimum of 4 h of milling is required to activate the solution mechanism in the Mg–6Zn system. However, 8 h of ball milling is necessary for the intermetallic phase to become visible. Upon closer examination of the main Mg peak (101), changes such as peak shifts, broadening, and intensity reduction are observed in the peak profile as the milling time increases (Figure 3.b). Peaks in the XRD data of pure elements typically appear at specific angles, reflecting homogeneous stresses in the lattice [38].
In the ball milling process, peak shifts occur primarily due to the increase in lattice stress, which is caused by mechanisms such as excessive plastic deformation, cold welding, or the formation of solid solutions from the dispersion of reinforcement elements within the primary structure [39]. For the Mg–6Zn system, the main peak shifts to lower angles as the milling time increases. This shift suggests that the distance between the planes in the lattice has widened [40,41]. The atomic structure diameter of Mg is 1.73 Å, while the atomic radius of Zn is 2.01 Å. These data suggest that the solid solution mechanism, which involves the dissolution of Zn into the Mg matrix, triggers the expansion of the unit cell [42]. The reduction in peak broadening and intensity is directly related to inhomogeneous residual stresses and crystallite size (CDDS).
The (002) peak of Mg in the XRD spectrum provides crucial information regarding the ball milling process theory (Figure 3c). Up to 4 h of milling, an increase in the intensity of the (002) peak is observed. This trend can be attributed to an increase in the crystal size of the hexagonal close-packed (HCP) structure of Mg. After 4 h of milling, SEM images reveal that the powder morphology transforms into a flaky or platelet structure due to excessive plastic deformation. In the HCP structure, the (002) peak represents the (c) direction, which is governed by van der Waals forces [43]. Consequently, atoms subjected to excessive plastic deformation preferentially align in the (c) direction due to crystallographic orientation, leading to the formation of crystallographic texture. This increase in intensity observed in the (002) peak is directly related to the formation of crystallographic texture. Beyond 4 h of milling, the flake morphology of the powders begins to break down, resulting in the disappearance of the crystallographic texture. Moreover, the decrease in peak intensity and significant peak broadening in the XRD data for powders milled for 8 h provides evidence of the elimination of this texture formation.

3.3. Effect of Variation Milling Time on Sintered Mg–6Zn Alloy System

3.3.1. Microstructure and Intermetallic Development in Sintered Materials

SEM-Mapping images of bulk samples obtained from pure Mg and ball milled powders at different milling times, following cold pressing and sintering processes, are shown in Figure 4. The SEM analysis of the pure Mg sample sintered for 2 h in an atmosphere-controlled tube furnace displayed the existence of pores, which are attributed to insufficient intragranular bonding. This lack of bonding appears to have led to intergranular separations. Comparable microstructural features were also reported in the study by Güneş et al., supporting the observations made in the current work [44]. According to the Mg–Zn phase diagram, 6.2 wt.% of Zn can dissolved in the Mg phase at 340 °C [45]. At room temperature, 2 wt.% of Zn can be found dissolved in the Mg phase. It can be stated that after the sintering process, as the temperature (500 °C) decreases to room temperature, the Zn atoms dissolved in the Mg phase rearrange to form MgZn intermetallics. In general, the internal structure of the Mg–6Zn alloy at room temperature consists of α-Mg and intermetallic phases. Upon examination of the SEM-Mapping images, it can be stated that the intermetallic structure formed in all three samples precipitates at the grain boundaries rather than within the grains.
The grain boundaries are regions with higher energy due to the presence of atom groups where bond formation is incomplete, making them more reactive. Therefore, the precipitates tend to form at these high-energy grain boundary regions in order to reduce the overall energy of the system [46,47]. Additionally, diffusion is more rapid at the grain boundaries [48]. Powder size and morphology before sintering, distribution of the reinforcement element, and the formation and distribution of the intermetallic phase are the primary factors influencing packability and sinterability.
High packability and successful sintering are crucial factors that determine the final microstructure. Furthermore, the final microstructure of the material is directly related to its mechanical, physical, and chemical properties [49].
The internal structure of the sample produced through sintering powders subjected to a 2 h ball milling process consists of Mg (gray regions) and intermetallic (white regions) phases, which precipitate along the grain boundaries. Furthermore, it can be stated pores are observed primarily at the grain boundaries. It is well known that the presence of these pores negatively affects the packability of the material. [50]. Based on the relevant image, it can be concluded that the quantity of pores at the grain boundaries decreases in the internal structure of the samples milled for 4 h and then sintered. Furthermore, larger grain sizes are observed, likely due to the flake or platelet structure of the powders after milling. Both the reduction in pore quantity and the coexistence of large and small grains indicate that the powders possess high packability. Materials sintered using powders milled for 8 h show a smaller grain formation, and the distribution of the Zn reinforcement element (lighter regions) is highly homogeneous. Additionally, it is evident that the secondary phase precipitated along the grain boundaries in these materials is smaller and more point-like compared to those in the other samples. Increasing the milling time to 8 h likely reduces the powder size of the reinforcement element as well as the matrix. This allows for better dispersion of the dissolved reinforcement element in the matrix during the re-deposition phase of liquid phase sintering. However, the packability of the powders is reduced due to excessive plastic deformation and work hardening [51]. Therefore, it can be said that a complete union between the particles is not achieved and voids occur in the structure, especially at the grain boundaries.

3.3.2. Effect of Milling Time on Hardness and Density in Bulk Materials

The hardness results for the samples belonging to the Mg–6Zn material system produced in this study are shown in Figure 5. From the graph, it can be observed that all hardness values increase with the milling time up to 8 h higher hardness values were achieved for all samples compared to pure Mg. As the milling time increases, the hardness values rise by 41.5% (55 ± 1.2HB), 84.6% (72 ± 1.65 HB), and 100% (78 ± 1.5 HB), respectively, compared to the pure form (39 ± 1 HB).
Such an increase in hardness can be attributed to the intermetallic phase formed in the structure after the ball milling process and sintering. Du et al. stated that there was an increase in hardness due to the formation of intermetallics [52]. Excessive plastic deformation resulting from repeated ball–powder–wall collisions may be one of the main factors contributing to the increase in hardness of the Mg–6Zn material system [53]. Due to the nature of the ball milling process, the powders are subjected to work hardening, where both hardness and brittleness increase. Additionally, the increase in hardness is directly related to the activation of dispersion hardening and solid solution hardening mechanisms. Lattice strain and dislocation density increase in the structure, especially after longer milling times during the ball milling process [19]. The Zn reinforcement element and intermetallic phases, which are uniformly distributed in the matrix due to the increasing milling time, create a barrier against dislocation movement. Different researchers have observed that in particle-reinforced metal matrix composite materials, the reinforcing element distributed homogeneously within the structure prevents dislocation movements [54,55]. This barrier leads to an increase in dislocation density around the precipitates, as described by the Orowan mechanism. The occurrence of concomitant strain hardening as a result of the Orowan mechanism has been reported in Mg–Zn systems by both Esteban et al. [56] and Alizadeh et al. [57]. Sübütay et al. [58] reported that the increase in dislocation density with prolonged ball milling time activated the grain refinement mechanism, leading to the formation of numerous sub-grains within the Mg matrix. The increasing dislocation density continues its path either by cutting through the particle or by forming a ring, depending on the size and strength of the precipitate. In both cases, the tension on the material must increase for the dislocation to progress [59]. Therefore, the hardness of the material also increases. Additionally, the smaller the powder size after milling, the smaller the grain size formation observed after sintering. According to the Hall–Petch equation, the relationship between grain size and hardness can be clearly demonstrated. The equation shows that a decrease in grain size results in an increase in hardness [60].
Figure 6 shows the relative density values of the samples from the Mg–6Zn system after the sintering process. In the powder metallurgy method, where both ball milling and sintering steps are used together, the mechanical and physical properties of the produced material system are highly dependent on the structural morphology of the powder and the ball milling parameters [61]. Properties such as powder morphology, the distribution of the reinforcement element within the structure, and its interaction with the matrix are the key parameters considered for the material. Process-based parameters such as ball milling time, lubricant type and amount, and ball-to-powder ratio have a direct effect on the microstructure of the material [62]. The relative density of pure Mg was measured to be 88%. This low density is likely due to the selection of a 500 °C sintering temperature and 2 h of sintering time. From this result, it can be concluded that both sintering temperature and time have a direct effect on the production of porous or dense materials. Prokopiev et al. [63] stated in their study that the material density can be controlled by adjusting the sintering temperature. It is preferred for materials used in the human body to be porous rather than having a high density [64]. For this reason, the corrosion rate values are considered more important than the density values obtained in the study.
In addition to temperature and time, the initial particle size and shape of pure Mg affect the density values of the final product. According to Doğan et al. [65], sintered samples composed of relatively larger particles were unable to reach full density. The density reached 94% in the sample subjected to 2 h of ball milling. In the ball milling process up to 4 h, the presence of platelet morphology powders and smaller particles broken from these pieces together in the structure can fill the gaps and improve packability. Accordingly, the density increased from 94% to 97%.
The platelet structure transforms into monodisperse spherical particles as the milling time increases. According to the literature, the initial powder size and distribution affect sinterability, packability, and density [66]. A structure with an unsuitable powder morphology will result in the formation of a weak interface between the matrix and the reinforcement element [67]. It is observed that the density decreases to 95.7% as the milling time increases to 8 h. This decreasing trend can be attributed to two factors. The first is the excessive plastic deformation that occurs after 8 h of milling, which reduces the compressibility of the powders due to cold welding. A similar effect was observed by Xu et al. in Mg/C nanofiber composites. They reported that the increase in milling time causes the powders to be cold welded and thus the compressibility to decrease, which leads to the formation of pores during sintering and a decrease in density [68]. As a result of this decrease in compressibility, the mechanical locking between the powders during pressing is weakened, causing the powders to repel each other and create gaps between them. These gaps persist as pores at the grain boundaries during sintering, thereby reducing the density. The second factor is that the intermetallic phase formed after 8 h of milling has a larger unit cell volume compared to the Mg and Zn phases. This may also contribute to the observed decrease in density.

3.3.3. Effect of Ball Milling Time on Degradation Behavior of Mg–6ZN Material System

The degradation products resulting from corrosion in the Mg–6Zn alloy system directly affect its biocompatibility properties. The degradation rate and behavior significantly influence mechanical stability and effective service life. Therefore, conducting research on corrosion products, rates, and mechanisms is essential. The immersion method used in this investigation is more meaningful as a reference compared to other methods [69]. The immersion test is widely preferred for biodegradable alloys due to its simple setup, ease of calculating average corrosion rates, and the ability to examine corrosion behavior over varying immersion times [70,71,72,73,74]. SEM images of the areas corroded on the surface of samples from the Mg–6Zn alloy system, which were immersed in Hank’s solution for 10 days, are shown in Figure 7. Upon examining these images, it can be observed that the white areas correspond to corrosion products. Cracks are particularly evident in the samples where powders ground for 2 and 4 h were used. It is believed that these cracks result from the H2 gas released during corrosion, a statement supported by various studies in the literature. [75,76,77]
The cracks that occur at the grain boundaries in samples produced from powders milled for 4 h appear thinner. Additionally, it is evident from the relevant image that pits form within the structure of this sample. In the sample subjected to 8 h of milling, it is observed that in addition to crack formation, separations between layers occur. It can be concluded that the intermetallic phase formed within the structure of this material contributes to corrosion by forming a galvanic couple with the Mg matrix. The graph showing the weight loss measurements of the samples produced in this study, presented as a percentage every 24 h, is shown in Figure 8a. The high weight loss indicates that the corrosion resistance of the material is low [78]. Based on this information, it can be argued that the 2 h ball milling process has a detrimental effect on corrosion resistance. As mentioned in previous sections, in the sample produced from powders milled for 2 h, the platelet morphology of the structure causes pores to form as a result of sintering (Figure 4). According to Wen et al. [79], the presence of these pores weakens the interfacial bonding between the matrix and the reinforcement phase. Consequently, these pores in the structure act as crack initiation sites during corrosion. Additionally, the high number of pores in the structure may lead to cracks merging and propagating along the surface. It can be concluded that such progression causes the structure to completely dissolve by the 10th day. When examining the sample produced from powders milled for 4 h, it can be stated that the corrosion resistance is higher, attributed to factors such as the robustness of the interface between the matrix and the reinforcement, which results from the most suitable packability in the structure, and the lower number of pores due to the presence of both large and small particles within the structure.
In the sample where the milling time was increased to 8 h, it was observed that the strength decreased rapidly after the 10th day. This decrease is thought to be due to the high amount of intermetallic phase formed in the structure. Additionally, the increased interaction between the Mg and Zn phases, due to cold-welded particles in the structure, may raise the probability of intermetallic phase formation. This, in turn, can lead to the formation of a more aggressive type of corrosion, specifically galvanic corrosion, by creating a galvanic couple. The graph showing the effect of milling time on the corrosion rate is presented in Figure 8b. The highest corrosion rate was observed in pure Mg. Ghali et al. reported that the corrosion behavior of pure magnesium and its alloys which is significantly affected by their microstructural characteristics, such as grain boundaries, phase distribution, and porosity [80]. This relationship between microstructure and corrosion resistance is further supported by the findings of Hook et al. [81], as well as the results obtained in the present study. Especially, Zhang at al. [82] stated that corrosion commonly initiates in the α-Mg phase, which acts as the anodic region in galvanic interactions with secondary phases. Hence, the higher corrosion rate observed in the pure Mg sample can be attributed to the dominance of the α-Mg phase within its microstructure. As stated in the previous section based on the SEM images (Figure 4), a significant number of pores were observed in the sample obtained by sintering the powders exposed to a 2 h ball milling process. Immersion tests revealed that the most rapid material degradation occurred by the 5th day for this sample (Figure 8a). Song et al. stated that pores concentrated along grain boundaries are particularly susceptible to attack by aggressive ions in the immersion environment, which would further compromise structural integrity [83]. These pores act as preferential sites for crack initiation and propagation. It can be stated that a significant amount of mass loss occurred in the sample due to these regions. After the 10th day, the mass loss rate tends to stabilize, which can be attributed to the detachment of highly porous regions and the formation of a corrosion product layer that limits further interaction with the medium. The results clearly indicate that the milling process, up to 4 h, positively enhances the corrosion resistance. However, as the milling time increases to 8 h, a noticeable increase in the corrosion rate of the Mg–6Zn alloy is observed. The occurrence of non-homogeneous corrosion in this sample can be attributed to the secondary phases precipitated along the grain boundaries [74]. Figure 9 shows schematically the surface corrosion between the matrix and the intermetallic phase.
These secondary phases act as cathodes, while the Mg matrix serves as the anode [84]. The secondary phases, acting as cathodes of galvanic corrosion, can form pits. Furthermore, the pores already present in the structure provide a favorable medium for the propagation of these pits [85]. As a result, in the sample produced from 8 h milled powder, the corrosion product layer is so loose that it cannot effectively protect the secondary phase matrix, leading to rapid dissolution [86]. By considering the results in this study, it can be inferred that the amount of porosity within the structure increases over time due to the effects of pitting corrosion. Similarly to the sample sintered for 2 h, the increasing porosity may serve as a site for crack initiation and propagation, ultimately compromising structural integrity and causing a sharp rise in mass loss. As a result, a sudden increase in mass loss is observed by the end of the 10th day.

4. Conclusions

In the presented study, Mg matrix materials reinforced with 6 wt% Zn at varying milling times were produced by means of the powder metallurgy method, which includes ball milling and sintering steps. Our aim is to determine the conditions under which optimum packability, hardness, density, and dissolution rate are achieved in this material system, where the reinforcement element is homogeneously distributed within the structure. The results presented below highlight the innovation of this study, focusing on the physical, chemical, and mechanical properties of the material system. During the milling process, platelet powder morphology is preserved for up to 4 h of ball milling due to the Mg phase’s ability to form texture in the (002) direction.
After a total of 8 h of milling time, the intermetallic phase along the grain boundaries within the structure is visible. This phase first appeared in the XRD results of powders milled for 8 h. In the case of the 2 h and 4 h ball milling process, it does not trigger this phase to become visible. The highest relative density value was obtained in the material produced under 4 h milling conditions (97.35%), where the reinforcement element was homogeneously distributed within the matrix. It can also be stated that optimum packability was achieved due to the presence of both small and large particles in the SEM images for this milling time. The highest hardness value was obtained in the material produced from powders milled for 8 h (78 ± 1.5HB) due to excessive plastic deformation and work hardening. In the corrosion rate and mass loss measurements, the most ideal material system was obtained in the 4 h ball milled samples (respectively, 6.51 mm/year and 28.47%).
By considering the results, these alloys not only enhance the hardness and density of pure Mg but also provide superior corrosion resistance, making them promising candidates for use in biomedical and clinical settings.

Author Contributions

S.B.Ç.: investigation, methodology, data curation, writing—original draft. E.S.: investigation, conceptualization, supervision, writing—review and editing. G.A.: investigation, methodology, data curation. A.D.: investigation, data curation. S.S.: investigation, data curation. H.S.: investigation, methodology, data curation, writing—original draft. All authors have read and agreed to the published version of the manuscript.

Funding

The financial support provided to this study by the Scientific Research Projects CoordinationUnit (SRPCU) of Selçuk University through contract# 22401147.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data will be made available on request.

Conflicts of Interest

The authors declare no conflict of interest.

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Coatings 15 00561 g001
Figure 1. Initial Powders (a) Mg, (b) Zn and ball milling powders in different milling time durations (c) 2h, (d) 4h and (e) 8h.
Figure 1. Initial Powders (a) Mg, (b) Zn and ball milling powders in different milling time durations (c) 2h, (d) 4h and (e) 8h.
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Coatings 15 00561 g002
Figure 2. SEM-Mapping images according to increasing milling time.
Figure 2. SEM-Mapping images according to increasing milling time.
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Figure 3. XRD representation (a) and peak profile exhibition of (101) (b) and (002) (c) for different milling durations.
Figure 3. XRD representation (a) and peak profile exhibition of (101) (b) and (002) (c) for different milling durations.
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Figure 4. SEM microstructure illustration of sintered Mg–6Zn alloy system.
Figure 4. SEM microstructure illustration of sintered Mg–6Zn alloy system.
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Figure 5. Hardness results of Mg–6Zn material system.
Figure 5. Hardness results of Mg–6Zn material system.
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Figure 6. Density results of Mg–6Zn material system.
Figure 6. Density results of Mg–6Zn material system.
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Figure 7. SEM images of surface morphology of Mg6Zn alloy system after 10 days of storage in Hank’s solution.
Figure 7. SEM images of surface morphology of Mg6Zn alloy system after 10 days of storage in Hank’s solution.
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Figure 8. (a) Weight % losses of Mg–6Zn alloys in Hank’s solution with time. (b) Corrosion rates of Mg–6Zn alloy samples.
Figure 8. (a) Weight % losses of Mg–6Zn alloys in Hank’s solution with time. (b) Corrosion rates of Mg–6Zn alloy samples.
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Figure 9. Schematic representation of surface corrosion depending on immersion time (a) 0 h, (b) 120 h, (c) 168 h, and (d) 240 h.
Figure 9. Schematic representation of surface corrosion depending on immersion time (a) 0 h, (b) 120 h, (c) 168 h, and (d) 240 h.
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Table 1. Ball milling parameters used in the study.
Table 1. Ball milling parameters used in the study.
Ball Milling Parameters
Milling Speed300 rpm
Ball/Powder ratio10:1
Chamber Fill10 g
PCA2 wt.% Stearic Acid
Ball Size10 mm
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MDPI and ACS Style

Çetinkal, S.B.; Salur, E.; Arıcı, G.; Degnah, A.; Sarkar, S.; Sübütay, H. Development of Zn-Reinforced Mg Matrix Composites via High Energy Ball Milling Duration: Impact on Mechanical Properties and Biodegradability. Coatings 2025, 15, 561. https://doi.org/10.3390/coatings15050561

AMA Style

Çetinkal SB, Salur E, Arıcı G, Degnah A, Sarkar S, Sübütay H. Development of Zn-Reinforced Mg Matrix Composites via High Energy Ball Milling Duration: Impact on Mechanical Properties and Biodegradability. Coatings. 2025; 15(5):561. https://doi.org/10.3390/coatings15050561

Chicago/Turabian Style

Çetinkal, S. Bilal, Emin Salur, Gökhan Arıcı, Ahmed Degnah, Sayan Sarkar, and Halit Sübütay. 2025. "Development of Zn-Reinforced Mg Matrix Composites via High Energy Ball Milling Duration: Impact on Mechanical Properties and Biodegradability" Coatings 15, no. 5: 561. https://doi.org/10.3390/coatings15050561

APA Style

Çetinkal, S. B., Salur, E., Arıcı, G., Degnah, A., Sarkar, S., & Sübütay, H. (2025). Development of Zn-Reinforced Mg Matrix Composites via High Energy Ball Milling Duration: Impact on Mechanical Properties and Biodegradability. Coatings, 15(5), 561. https://doi.org/10.3390/coatings15050561

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