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Article

Study on the Effect of Pressure on the Microstructure, Mechanical Properties, and Impact Wear Behavior of Mn-Cr-Ni-Mo Alloyed Steel Fabricated by Squeeze Casting

by
Bo Qiu
*,
Longxia Jia
,
Heng Yang
,
Zhuoyu Guo
,
Chuyun Jiang
,
Shuting Li
and
Biao Sun
*
School of Mechatronic Engineering, Xi’an Technological University, Xi’an 710021, China
*
Authors to whom correspondence should be addressed.
Metals 2024, 14(9), 1054; https://doi.org/10.3390/met14091054
Submission received: 21 August 2024 / Revised: 7 September 2024 / Accepted: 12 September 2024 / Published: 15 September 2024
Browse Figures
Figure 1
<p>Sampling position diagram of the prepared sample.</p> "> Figure 2
<p>Microstructure of the samples prepared under different pressures: (<b>a</b>) 0 MPa; (<b>b</b>) 30 MPa; (<b>c</b>) 60 MPa; (<b>d</b>) 90 MPa; (<b>e</b>) 120 MPa; (<b>f</b>) 150 MPa.</p> "> Figure 3
<p>Secondary dendrite arm spacing and ferritic content of the samples prepared under different pressures.</p> "> Figure 4
<p>SEM images at a higher magnification of the microstructure of the samples prepared under different pressures: (<b>a</b>) 0 MPa; (<b>b</b>) 30 MPa; (<b>c</b>) 60 MPa; (<b>d</b>) 90 MPa; (<b>e</b>) 120 MPa; (<b>f</b>) 150 MPa.</p> "> Figure 5
<p>XRD analysis of the steel prepared under various pressures.</p> "> Figure 6
<p>Variation in the density and porosity of the steels prepared at different pressures.</p> "> Figure 7
<p>(<b>a</b>) Brinell hardness of the samples prepared under different pressures; (<b>b</b>) low-temperature (−40 °C) impact energy of the samples prepared under different pressures.</p> "> Figure 8
<p>Macro- and micro-impact fracture morphology of the samples prepared under different pressures: (<b>a</b>–<b>a’</b>) 0 MPa; (<b>b</b>–<b>b’</b>) 30 MPa; (<b>c</b>–<b>c’</b>) 60 MPa; (<b>d</b>–<b>d’</b>) 90 MPa; (<b>e</b>–<b>e’</b>) 120 MPa; (<b>f</b>–<b>f’</b>) 150 MPa.</p> "> Figure 9
<p>EDS analysis results of the impact fracture morphology for the sample prepared under 30 MPa: (<b>a</b>) site 1; (<b>b</b>) site 2.</p> "> Figure 10
<p>(<b>a</b>) Relationship between wear time and wear loss of the samples prepared under different pressures; (<b>b</b>) wear rate of the samples prepared under different pressures.</p> "> Figure 11
<p>Morphology of the worn surface of the samples prepared under different pressures: (<b>a</b>) 0 MPa; (<b>b</b>) 30 MPa; (<b>c</b>) 60 MPa; (<b>d</b>) 90 MPa; (<b>e</b>) 120 MPa; (<b>f</b>) 150 MPa.</p> ">
Versions Notes

Abstract

:
ZG25MnCrNiMo steel samples were prepared by squeeze casting under pressure ranging from 0 to 150 MPa. The effects of pressure on the microstructure, low-temperature toughness, hardness, and impact wear performance of the prepared steels were experimentally investigated. The experimental results indicated that the samples fabricated under pressure exhibited finer grains and a significant ferrite content compared to those produced without pressure. Furthermore, the secondary dendrite arm spacing of the sample produced at 150 MPa decreased by 45.3%, and the ferrite content increased by 39.1% in comparison to the unpressurized sample. The low-temperature impact toughness of the steel at −40 °C initially increased and then decreased as the pressure varied from 0 MPa to 150 MPa. And the toughness achieved an optimal value at a pressure of 30 MPa, which was 65.4% greater than that of gravity casting (0 MPa), while the hardness decreased by only 6.17%. With a further increase in pressure, the impact work decreased linearly while the hardness increased slightly. Impact fracture analysis revealed that the fracture of the steel produced without pressure exhibited a quasi-cleavage morphology. The samples prepared by squeeze casting under 30 MPa still exhibited a large number of fine dimples even at −40 °C, indicative of ductile fracture. In addition, the impact wear performance of the steels displayed a trend of initially decreasing and subsequently increasing across the pressure range of 0–150 MPa. The wear resistance of samples prepared without pressure and at 30 MPa was superior to that at 60 MPa, and the wear resistance deteriorated when the pressure increased to 60 MPa, after which it exhibited an upward trend as the pressure continued to rise. The wear mechanisms of the samples predominantly consisted of impact wear, adhesive wear, and minimal abrasive wear, along with notable occurrences of plastic removal, furrows, and spalling.

1. Introduction

Railway transport is often regarded as a preferred mode of transportation due to its developing industrialization, enhanced safety compared to other modes, lower transport cost, and higher transport efficiency [1]. As primary considerations in railway transport, safety and reliability are constant themes. Consequently, ensuring the reliability of railway systems and preventing the failure of structural components have consistently garnered significant interest [2,3,4,5]. As a crucial component that connects the carriages, the coupling experiences multi-directional dynamic loads, typically transmitted through the draw hook, with the hook tongue identified as its weak link. Therefore, enhancing the performance of the hook tongue has emerged as a significant concern in the field of railway transport. Chunduru et al. [6] concluded from their study that the actual service life of the hooks often falls short of expectations due to stress concentrations. They identified the hook tongue as the primary failure area, which showed significant improvements following enhancements made to its design. Consequently, enhancing the performance of the hook tongue, identified as a weak link, is advantageous for the overall performance of the hook. Huang et al. [7] experimentally concluded that the solidification microstructure of the failed hook tongue exhibited coarse dendritic formation and contained numerous shrinkage holes. The fracture mode was identified as brittle fracture. It is inferred that the fracture mechanism of the hook tongue involves micro-cracks generated during welding repair, which propagate under impact loads along the shrinkage holes, ultimately leading to instantaneous failure. Therefore, there is an urgent necessity to enhance the comprehensive mechanical properties of the hook tongue while ensuring its strength and wear resistance, thereby achieving high strength alongside improved plasticity, low-temperature impact toughness, and fatigue resistance. Several researchers have indicated that the improper execution of regular maintenance can lead to fatigue damage or even failure of the hook tongue, thereby putting transport safety at risk [8,9].
ZG25MnCrNiMo is the primary material used for the hook tongue of rail vehicles, exhibiting characteristics such as high strength, favorable forming properties for both hot and cold machining, good corrosion resistance, and high fatigue strength [10]. The development of the railway industry toward high-speed and heavy-duty operations has significantly deteriorated the service conditions of such material. Low temperatures, heavy loads, and prolonged exposure to high-intensity and alternating loads have led to failures in the hooks, including breaking and other serious accidents. These issues expose some drawbacks, including insufficient toughness, poor impact property stability, and weak resistance to low-temperature cracking [11], which pose risks to railway transportation safety. Therefore, enhancing the comprehensive mechanical properties of ZG25MnCrNiMo steel, particularly its toughness and impact resistance, holds significant practical importance for achieving substantial advancements in railway transport. Currently, extensive research has been conducted on the composition enhancement, microstructural optimization, mechanical property improvement, and production process optimization of ZG25MnCrNiMo steel used in rail transportation. Wang [12] found that the inclusions in such steel for railway wagons are primarily non-metallic inclusions of oxides and silicates, exhibiting various forms and sizes with a uniform distribution. Although they provided explanations for the causes of these inclusions and corresponding control measures, fundamental solutions to the casting inclusion problem remain unaddressed. Zhang [13] investigated the −40 °C impact toughness of ZG25MnCrNiMo steel for hooks subjected to various tempering processes, and the results indicated that optimal impact performance was achieved when the material was quenched at 910 °C followed by tempering at 550 °C, resulting in the formation of a tempered sorbite microstructure, with impact work increasing alongside tempering temperatures. Despite the improvements in toughness for this steel, the effects are not pronounced. Therefore, to enhance toughness further, many researchers have adopted forging methods to prepare components. Zhang [14] assessed the fatigue strength of this steel produced through both forging and casting processes, determining fatigue life under these two conditions and performing S-N curve fitting. The results indicated that the forged steel exhibited superior fatigue performance compared to the casted steel. However, the forging process involves high costs and a complicated procedure for complex hook tongues, making it challenging to promote forged hook tongues in the market.
The rail industry has been actively engaged in the development of innovative preparation technology aimed at enhancing the service life and reliability of hook tongues to ensure the safety of rail transport. In current study, the squeeze-casting process was proposed to fabricate ZG25MnCrNiMo steel parts, providing a reference for enhancing the microstructure and overall mechanical properties of complex components in process development. Squeeze casting is a metal-processing technique wherein a volume of molten metal is directly injected into a high-strength pressure chamber or mold cavity. Subsequently, mechanical static pressure is continuously applied, allowing the molten metal to fill the mold and solidify. This process compensates for shrinkage and induces minimal plastic deformation due to exerted pressure, resulting in components with a dense microstructure, smooth and clean surfaces, and precise dimensions [15]. Additionally, it offers several advantages, including high productivity, reduced casting defects, excellent product performance, high material utilization, and low production costs. Numerous studies have demonstrated that squeeze casting can substantially enhance the strength and toughness of non-ferrous metals and their composite materials [16,17,18]. Zhao [19] studied the microstructure and mechanical behavior of an AZ91D alloy under different applied pressures by near-liquidus squeeze casting, and he found that the ultimate tensile strength (UTS) and elongation (EL) of differential support can be improved by controlling the real-time precision of the forming pressure. The UTS and EL of the alloy reached the maximum values when the applied pressure was 140 MPa, and the comprehensive mechanical properties were the best. For ferrous metals, however, high melting points and poor fluidity pose significant challenges for squeeze casting, leading to slow advancements in the research and application of this process for ferrous metals. Additionally, the influence of applied pressure on the microstructure and mechanical properties of cast steel materials, particularly regarding low-temperature impact toughness, has not yet been documented.
In the current study, the squeeze-casting process was employed to manufacture ZG25MnCrNiMo steel, taking into account its low costs and high-forming-quality characteristics. Subsequently, the crucial controllable process parameter, squeezing pressure, and its influence on the microstructure, low-temperature impact toughness at −40 °C, and wear properties of ZG25MnCrNiMo steel were investigated, thereby establishing a foundation for the trial production of ZG25MnCrNiMo steel hook tongues using the squeeze-casting method.

2. Materials and Methods

2.1. Materials and Experiment

A KGPS 25 kg medium-frequency induction furnace (Zhengxin, Dongguan, China) was used to melt the alloy materials in an air environment. Initially, a round bar of ZG25MnCrNiMo steel material was placed into a corundum crucible. Following melting and skimming, based on the quantity of alloying elements consumed, ferromolybdenum, ferrochrome, electrolytic nickel plate, and ferromanganese were added as replenishments, alongside a carbon additive to compensate for the burned carbon. The molten steel was poured at a temperature of 1650 °C, after which the steel was transferred to a ladle and subsequently poured into a metal mold. And the ladle and mold were preheated to 700 °C and 200 °C, respectively. Then, a THP16-200A four-column vertical squeeze-casting machine (Tianduan, Tianjin, China) was employed to conduct the direct squeeze casting of the molten metal, and the sample was ejected following solidification. The exerted pressure during squeeze casting was set as 0 MPa, 30 MPa, 60 MPa, 90 MPa, 120 MPa, and 150 MPa, and the holding time of pressure was maintained at 30 s. And for different prepared conditions, the same furnace molten steel and mold were employed for casting. The Spectro-MAXX direct-reading spectrometer (Spectro, Klevel, Germany) was utilized to analyze the chemical composition of the prepared steel, which is presented in Table 1.

2.2. Instruments and Characterizations

The metallographic samples, low-temperature impact toughness samples, and wear samples were cut using a wire cutting machine from the center contact area with the pressure head, as depicted in Figure 1. Metallographic samples with size of 10 mm × 10 mm × 30 mm were prepared and etched with 4% Nital solution, followed by polishing. Then, the microstructure of samples was observed using a Leica DM2000X optical microscope (OM, Leica, Weztlar, Germany) and a Zeiss EVO18 scanning electron microscope (SEM, Carl Zeiss, Heidenheim, Germany) equipped with an energy dispersive spectrometer (EDS, Bruker Nano GmbH, Berlin, Germany). The quantitative metallographic analysis was based on GB/T 15749–2008 [20], and the images were imported into Image-Pro-Plus software (6.0, Media Cybernetics company, Rockville, MD, USA) for acquisition, analysis, and quantitative calculations. To improve the accuracy of the metallographic measurement results, each sample was photographed with 15 fields of view, and the measurements included (a) measurement of the secondary dendrite arm spacing (SDAS) by the truncation method and (b) measurement of the ferrite content by the area method. In addition, the phases of the samples were determined using an Ultima IV XRD (Rigaku, Tokyo, Japan) with the following parameters: Cu Kα radiation, a voltage of 40 kV, a current of 200 mA, a scanning angle range of 40° to 140°, and a scanning speed of 10°/min.
According to the Archimedes drainage method for measuring the density of the samples prepared under different pressures, a precision electronic balance was utilized to weigh the sample’s weight; subsequently, the following formula was applied to calculate the actual density of the samples:
ρ 1 = ρ 2 · m 1 m 1   − m 2
where ρ 1 represents the actual density of the sample (g/cm3); ρ 2 denotes the density of water (g/cm3); and m 1 and m 2 signify the mass of the sample in water and air, respectively. Further, the formula for calculating the porosity of prepared samples is as follows:
f = ρ t − ρ 1 ρ t
where f signifies the porosity (%); ρ t represents the theoretical density of the sample (g/cm3). And the theoretical density can be calculated based on the chemical composition of the material and the densities of its constituent elements. After calculating, the theoretical density of the steel was determined to be 6.8424 g/cm3. In addition, three samples were taken from steel prepared under each pressure to measure the density and porosity, and then the average values were reported as the results.
Hardness tests were conducted using the metallographic samples, and Brinell hardness was measured using a THBRV-187.5 electrically operated Brodmann hardness tester (Hengyi, Beijing, China) for the samples prepared under various pressures. The indenter is made of 2.5 mm diameter tungsten carbide balls, with a test force of 187.5 kg and a holding time of 25 s. Ten points were measured for each sample, and the average was taken as the final result. The low-temperature impact toughness samples were machined into a 10 mm × 10 mm × 55 mm standard size with a V-notch depth of 2 mm. The samples were cooled in a mixture of ethanol and liquid nitrogen at −40 °C and held for 10 min to make the samples cool evenly; then, the experiment was conducted in a JBDS-300B low-temperature impact tester (Zhongnuo Instrumentation Co., Ltd., Jingjiang, China). After the impact test, macroscopic fractures were observed with an S6D stereomicroscope, and microscopic fractures was observed with a Zeiss EVO48 SEM (Carl Zeiss, Heidenheim, Germany) and analyzed by EDS (Bruker Nano GmbH, Berlin, Germany).
Impact wear is a combined action of impact and slide wear, which is closer to the actual working conditions of the draw hook. In current study, the impact wear test was conducted on an MLD-10 dynamic load abrasive wear tester (Chengxin Testing Equipment Co., Ltd., Zhangjiakou, China) to investigate the wear behavior of the prepared samples. During the wear process, the upper sample (tested materials with a size of 10 mm × 10 mm × 30 mm) reciprocated with the hammer, while the lower sample (counterpart ring) rotated with the shaft to contact the lower part of the upper sample. And the counterpart ring with a size of 50 mm in outer diameter, 30 mm in inner diameter, and 20 mm in thickness was made from GCr15 steel with a hardness of HRC 62–64. Given that the actual working condition of the draw hook is contact wear without abrasive intervention, an abrasive-free impact wear test was carried out. According to the service condition of the hook, Equation (3) was used to calculate its relative impact work E:
E = ( F · x ) / S
where F represents impact force (N); x signifies the distance of action, here referring the clearance of the hook (mm); and S denotes the action area (mm2). And the impact force on the hook is about 1200 KN, the clearance between the hooks is about 6 mm, and the contact area is about 2800 mm2, so the calculated impact work E is 2.57 J/mm2. In the current experiment, the impact work was set as 2.5 J, which corresponded to free-fall heights of the impact hammer from 25 mm.
Consequently, the parameters for the impact wear test included the following conditions: an impact work of 2.5 J, an impact frequency of 150 times/min for the impact hammer, a rotation speed of 150 r/min for the counterpart ring, and no external abrasion. The parameters of the wear test are presented in Table 2. To ensure that the contact surface between the upper and lower samples transformed from line contact to face contact, 60 min of pre-wear was conducted prior to the test. After the pre-wear process, the samples were cleaned with ultrasonic treatment and then dried and weighed using an electronic balance with 0.1 mg precision before and after each test to measure the mass variation. The wear total time for each sample was 300 min, and the wear weight loss was measured at 20 min intervals. The average of the weight loss of the three wear samples at each pressure was used to assess the impact wear resistance of the material.

3. Results and Discussion

3.1. Microstructure

The microstructure of the tested materials under varying pressures was examined, as illustrated in Figure 2. Given that ZG25MnCrNiMo steel is classified as hypoeutectoid steel, it is characterized by a microstructure that comprises gray pearlite and a white network of ferrite. When no pressure is applied during preparation, the steel exhibits coarse dendritic grains (Figure 2a). With increasing pressure, the distance between the secondary dendrite arms decreases, resulting in a finer structure. The pressure causes the eutectoid point of the Fe-Fe3C system to shift toward a lower temperature and lower carbon content, resulting in a diminished α phase region [21]. Consequently, the carbon content of the prepared steel approaches eutectoid composition, facilitating the formation of a network ferrite that precipitates along the protoaustenite grain boundaries. As illustrated in Figure 2b–f, the overall uniformity of the solidification structure of the prepared steel under pressure demonstrates notable consistency without discernible directionality. This phenomenon can be attributed to the fact that crystallization under pressure minimizes the diffusion time of solutes in liquid steel [22], significantly reduces dendrite segregation, and enhances the uniformity of chemical composition. Moreover, given that the elongated dendrites are fused and remelted under the action of pressure, the directivity of the grains is diminished, thereby further promoting the homogenization of the microstructure.
The secondary dendrite arm spacing of prepared steel under varying pressures was quantitatively measured and statistically analyzed, and each sample was photographed taking 15 fields of view; the results are shown in Figure 3. In comparison to the sample prepared without applied pressure, the secondary dendrite arm spacing of the sample subjected to 30 MPa exhibits a decrease of approximately 32.4%. As the pressure continues to increase, the grain refinement effect becomes increasingly pronounced. And when the pressure is elevated to 150 MPa, the secondary dendrite arm spacing constitutes only 54.6% of that observed in gravity casting (0 MPa). This can be attributed to the applied pressure diminishing the nucleation activation energy [23], elevating the melting point of the alloy, and enhancing the degree of supercooling, thereby increasing the nucleation rate. In addition, the pressure augments the cooling rate, preventing the grains from growing adequately, thereby facilitating grain refinement [24].
The ferritic content of prepared steel under different pressures was also quantitatively determined (Figure 3). As the pressure increases, the ferrite content within the microstructure gradually increases, as reflected in the statistical results presented in Figure 3. Compared with that of the sample prepared without pressure, the ferrite content of the sample with a pressure of 150 MPa is elevated by 39.1%. The refinement of the original austenite grains results in an increase in grain boundaries, thereby facilitating the nucleation of ferrite in more locations. In addition, when the degree of supercooling is greater, the phase transition kinetics of the austenite–ferrite transformation depend on the rate of interfacial reaction [25]. The pressure induces thermal deformation of the austenite phase and enhances its chemical potential. The increment in chemical potential, Δd, can be expressed as shown in Equation (4) [26]:
∆ d = G d b 2 V V / 2
where VV represents the molar volume of austenite phase; G and b denote the shear modulus and Bertrand vector of the austenite phase, respectively; and d indicates the dislocation density, which is contingent upon the deformation conditions. The work of deformation induced by pressure on the austenite is partially transformed into stored deformation energy. This transformation leads to an increase in dislocation density within the austenite phase and elevates its chemical potential. Consequently, the driving force for dynamic phase transformation is enhanced, which manifests as an increase in the amount of ferrite precipitation within the microstructure.
Figure 4 presents the SEM images of the microstructure of the prepared steels at a higher magnification. The prominent black network structure depicted in Figure 4 represents ferrite, while the black and white regions correspond to flake pearlite. As shown in Figure 4a, ferrite nucleates at the grain boundary of the protoaustenite and grows in a “block” form, resulting in a spun crystal morphology. With the increasing pressure, as depicted in Figure 4b–f, the isometric ferrite located at the grain boundary progressively transforms into a lamellar structure and intragranularly grows along the grain boundaries. Meanwhile, the lamellar spacing gradually diminishes, subsequently transforming into an acicular shape that integrates into the grain. This phenomenon results from a substantial increase in the degree of supercooling induced by the elevated pressure. A small quantity of proeutectoid ferrite is expelled from the austenite and precipitates along the grain boundary [27], forming network ferrite and acicular ferrite. Upon cooling below Ar1, eutectoid decomposition occurs, and the structure is transformed into a lamellar perlite structure. The integrated structure comprising intragranular acicular ferrite and extremely fine lamellar pearlite is referred to as the Widmannstatten structure [28]. Considering that the Widmannstatten structure can only be formed within a certain range of cooling rates, it requires a sufficiently high degree of supercooling to provide an adequate driving force for the phase transition, and then the coherent strain energy necessary to create the interface of the pin ferrite semi-common format becomes attainable [29]. Consequently, the accelerated cooling rate facilitated by pressure promotes the formation of the Widmannstatten structure, and with the increase in supercooling degree, this microstructure is progressively refined.
The XRD analysis of the samples prepared at various pressures is illustrated in Figure 5. With the increase in applied pressure, the broadening of the X-ray Bragg diffraction peaks of the sample become more pronounced, and the intensity of the diffraction peaks gradually increases, concurrent with the observed shift of the diffraction angle toward a lower value, as illustrated by the magnified α-Fe (110) peaks in the figure. The broadening of Bragg diffraction peaks is primarily influenced by grain refinement [30]; increased pressure promotes grain size refinement, contributing to the strengthening effect. Additionally, changes in diffraction intensity are influenced by various factors, including crystal structure, grain size, and dislocation density. As illustrated by the localized magnification of the diffraction spectra, the diffraction intensity of the α-Fe (110) grain surface is significantly enhanced, indicating that grain refinement increases the number of grain boundaries, facilitating the scattering of X-rays by atoms in the crystals and thereby enhancing the diffraction intensities.

3.2. Mechanical Properties

The density of the alloy serves as a crucial indicator of the porosity of the casting. The variation in the density and porosity of the prepared steel under different pressures is illustrated in Figure 6. It is evident that as the pressure increases, the density of the steel rises, while conversely, the porosity of the castings decreases. During the squeeze-casting process, the application of pressure not only increases the density of the casting but also reduces the average distance between atoms, thereby minimizing macro defects in the casting. Under a pressure of 150 MPa, the sample’s density attains a maximum value of 6.8135 g/cm3, while its porosity achieves a minimum of 1.33%, indicating that the casting exhibits optimal densification. Meanwhile, the density and porosity of the steel exhibit the most significant changes when the pressure is increased from 0 to 30 MPa; the density increases by 1.61%, while the porosity decreases by 49.5%. Furthermore, the rate of change in density and porosity diminishes as the pressure continues to rise.
There are several factors contributing to porosity, primarily including two main reasons: First is the gas dissolved in the alloy. During the melting process of the alloy, the liquid metal contains a higher concentration of dissolved gas, which is correlated with the temperature of the molten metal. As the temperature of the metal rises, the solubility of the gas increases. Additionally, during the pouring process, the gas tends to break down into numerous smaller monomers, which possess lower buoyancy and struggle to escape from the liquid surface, resulting in the formation of pores during the casting process. When the pressure is 0 MPa, the spacing between dendrites poses challenges for filling, making it difficult to eliminate shrinkage. However, when pressure is applied, the flowing distance for the molten metal is reduced. Consequently, less energy is consumed, facilitating easier filling of the liquid. In addition, pressure inhibits the formation of bubbles and increases the solubility of gas in the molten metal, thereby reducing the occurrence of defects such as shrinkage.
Figure 7a illustrates the Brinell hardness of prepared steel at various pressures. All samples exhibit hardness values within the range of HB241 to HB311. As the pressure increases, the hardness of the sample initially decreases before subsequently increasing. The hardness reaches a peak value of HBW286.4 at an applied pressure of 120 MPa, and as the pressure continues to increase, the hardness exhibits a slight decrease.
Figure 7b illustrates the influence of pressure on the low-temperature impact energy of the prepared steels. When the pressure is set at 30 MPa, the low-temperature (−40 °C) impact toughness of the steel reaches its maximum value of 31.79 J, representing an increase of 65.4% compared to the sample that exhibited a value of 19.22 J without pressure applied. The low-temperature impact toughness of the steels exhibits a linear decrease with increasing pressure. Consequently, the empirical relationship illustrated in Equation (5) can be derived through linear regression:
A k v = − 0.147 P + 36.933
where Akv represents low-temperature impact absorption energy (J); P denotes prepared pressure (MPa). The fitting slope of the empirical relationship is −0.147, with a correlation coefficient reaching 96.14%.
When the pressure reaches 60 MPa, the low-temperature impact absorption energy decreases to 27.21 J, approaching the lower limit of 27 J for ZG25MnCrNiMo steel, as stipulated for casting samples or Kiel test blocks in TB/T2942–1999 [31]. As the exerted pressure continues to rise, the low-temperature impact performance fails to satisfy the necessary requirements for application. In the current study, the maximum pressure for squeeze casting has been raised to 150 MPa, while the impact absorption energy remains merely 14.54 J. It should be noted that the impact toughness of samples is generally much lower than that of cast samples and Kiel test blocks, even when utilizing the same liquid metal [7]. In summary, it can be concluded that the low-temperature impact toughness of the test steel is significantly influenced by the applied pressure within the range of 0 to 150 MPa, and the impact property is superior under lower pressure. In addition, the impact toughness of the steel without applied pressure is also relatively low, failing to meet the required standards.
The macro-impact fracture of the tested steel at −40 °C was examined, as illustrated in Figure 8a–f. The fracture is characterized by obvious plastic deformation. The fiber region and shear lip of the fracture are formed following plastic deformation, and the two constitute the primary components responsible for impact energy absorption. In comparison to the schematic diagram of the impact fracture surface [32], it is evident that when the pressure ranges from 30 to 90 MPa (Figure 8b–d), the proportion of the fiber zone and shear lip in the impact fracture is larger, indicating relatively good toughness. As depicted in Figure 8e,f, the fiber area and shear lip of the macro fracture are diminished, resulting in relatively poor impact toughness under conditions of elevated pressure and without pressure applied (Figure 8a).
Additionally, Figure 8a’–f’ illustrates the impact fracture morphology at a higher magnification. The impact fracture of the sample prepared without pressure exhibits quasi-cleavage morphology, which is characteristic of a brittle transgranular fracture. Impurity particles located at the bottom of the secondary dendrite arms are distinctly observable, which is attributed to the concentration of solute elements in the liquid phase at the crystallization front, leading to dendrite segregation. At an applied pressure of 30 MPa, as depicted in Figure 8b’, the low-temperature impact toughness of the steel reaches the optimal value, and the fracture consists of numerous finer dimples, indicative of a typical ductile fracture. The increase in pressure shortens the solidification time, consequently decreasing the diffusion time of solute elements within the liquid phase, minimizing alloy segregation, and promoting the uniform distribution of solute elements. Consequently, the low-temperature impact toughness of the steel prepared under lower pressure significantly exceeds that of the steel prepared by gravity casting (0 MPa). However, as the pressure continues to increase, the dimples progressively become shallower, as shown in Figure 8c’. In Figure 8d’,e’, the material exhibits quasi-cleavage fracture characteristics, but notable plastic deformation remains observable at the edge of the tear ridge. And in Figure 8e’, cleavage steps, tongue-shaped patterns, tear ridges, and small cleavage faces are visible. When the pressure increases to 150 MPa, as depicted in Figure 8f’, numerous micro-cracks emerge at the fracture. This phenomenon occurs because the increase in pressure elevates the carbide content in the material and alters the carbide morphology, resulting in the quasi-cleavage cracks, which are more likely to originate from fine carbides within the crystal structure [33], thereby diminishing the toughness of the material.
EDS spectrum analysis was conducted on the particulate matter at the discontinuity in Figure 8b’, and the results are presented in Figure 9. The particles at site 1 predominantly consist of sulfide MnS and FeS. Given that the S readily segregates at the grain boundaries [34], the interfacial energy diminishes, consequently increasing the likelihood of crack propagation along the cleavage plane. The particles at site 2 are alloy cementite (Fe,Mn)3C with high hardness. Mn possesses a greater affinity for C than Fe and exhibits a higher propensity for segregation, facilitating precipitation along the grain boundary and forming the network carbides such as Mn3C or (Fe,Mn)3C. The precipitation of these particles adversely affects the low-temperature impact toughness of the tested steel to some degree, which can be mitigated by precisely controlling the chemical composition during the liquefaction process preceding the furnace.
During the squeeze-casting process, when the pressure is increased from 0 to 30 MPa, it can enhance the solidification rate of metals, increase the degree of supercooling, and facilitate the grain refinement of the steel. Meanwhile, this process reduces the stress concentration, increases the grain boundary ratio, and renders the crack growth path more tortuous and directed, thereby absorbing more impact energy and resulting in fine crystalline strengthening and toughening. Furthermore, the application of pressure induces the thermal deformation of the austenite, shifting the austenite–ferrite transformation toward increased ferrite formation. Consequently, this leads to an increased ferrite content, which enhances the impact toughness of the steel [35,36]. The pressure toughening effect is substantiated by the experimental results obtained in this study. When the pressure of squeeze casting is raised to 60 MPa, a minor quantity of the Widmannstatten structure precipitates along the grain boundaries and grows into the intragranular region, creating a partitioning effect on the grains and leading to the formation of cleavage fracture channels. The inherent morphology of the crystalline structure cannot be eliminated, resulting in a substantial reduction in the impact toughness of the prepared steel. The carbides with pointed features are prone to serving as crack initiation sites during impact, resulting in the fracture of samples. In addition, the pressure of squeeze casting intensifies the solid solution strengthening of the alloy, resulting in decreased plasticity and toughness [37]. Even at a pressure of 150 MPa, micro-cracks may develop, indicating pressure embrittlement. In summary, during the squeeze-casting process under lower pressure, fine crystalline strengthening and toughening predominantly influence the impact process, resulting in pressure toughening of the steel. Conversely, during high-pressure squeeze casting, solid solution strengthening and the appearance of the Widmannstatten structure play a crucial role, intensifying the embrittlement trend of the material and making the material behave according to pressure embrittlement.

3.3. Wear Property

The influence of wear time on the wear loss of the samples prepared under different pressures is plotted in Figure 10a. The wear resistance of the samples prepared under 0 MPa and 30 MPa is moderate. Excluding the pre-grinding process during the initial 60 min, the samples remain in a stable wear stage within 60 to 220 min, characterized by a low wear rate and a gradual increase in wear amount. After 220 min, the wear rate increases, leading to accelerated wear. When the pressure ranges from 60 to 90 MPa, the wear resistance of the tested steel is poor, characterized by a short stable wear stage and significant total wear. And when the pressure exceeds 120 MPa, the wear resistance of the tested steel is optimal. Even after 300 min of wear, the samples remain in a gentle stable wear stage, with total wear loss remaining below 50 mg.
The wear rate per unit area of the steel corresponding to pressures ranging from 0 to 150 MPa are illustrated in Figure 10b. The wear rate initially increases before subsequently decreasing as the pressure increases. When the pressure is increased within the range of 0 to 60 MPa, the wear rate increases by 151.6%, reaching its lowest wear resistance at 60 MPa, where the wear rate attains 40.25 mg/(h·cm2). When the pressure falls within the range of 60 to 150 MPa, the wear rate decreases by 83.2% as the pressure increases, ultimately decreasing to 6.75 mg/(h·cm2) at the high pressure of 150 MPa, indicating optimal wear resistance.
The worn surface morphology of samples after 300 min of impact abrasive wear under conditions of 2.5 J impact energy and various pressures was examined, as shown in Figure 11. Under the action of impact, the surface of the material primarily exhibits plastic deformation, and distinct extrusion marks are evident. Since the rotational direction of the lower sample in the experiment remains consistent, the extrusion marks display pronounced directional characteristics. As illustrated in Figure 11a,b, the worn surface of samples prepared without pressure and under 30 MPa squeeze casting is relatively flat, with wear debris that can push the material to one side, thereby forming a plastic ridge and wedge. A limited number of spalling pits resulting from matrix fractures and detachment due to insufficient strength are observable [38], and the furrows created by the scratches of the wear debris are shallow, indicating relatively good wear resistance. When the pressure reaches to 60 MPa and 90 MPa, as depicted in Figure 11c,d, the material experiences damage from stripping under impact, and the spalled material is rolled between contact surfaces, forming abrasive debris that exhibits significant scratching effects, causing deeper grooves, increased inclusions, and reduced wear resistance. When the pressure exceeds 120 MPa, as illustrated in Figure 11e,f, the wear surface becomes notably flat, with minimal damage caused by abrasive debris. The resulting minor deformation makes it challenging to remove the surface material, thereby leading to reduced wear loss.
The lower sample of the impact wear experiment is a 20Cr ring, which is also a kind of alloy steel material as the tested samples. Under the action of impact, significant local stress at the contact surface between the two materials is generated, resulting in a pronounced tendency for cold welding. Subsequently, the adhesion points fracture due to shear effects during relative sliding, leading to the formation of wear debris that further scratches the surface [39]. Then, the surface experiences a combined action of normal and tangential forces, resulting in a significant material loss and forming furrows and a limited number of inclusions. The wear mechanisms observed on the worn surface under the condition of 2.5 J impact energy primarily include impact wear, adhesive wear, and minimal abrasive wear, along with notable occurrences of plastic removal, furrows, and spalling.
When the pressure of squeeze casting is lower (38 MPa), the wear resistance of prepared steel does not significantly differ from that of gravity casting (0 MPa). It can be considered that low-pressure squeeze casting has a minimal effect on the wear resistance of ZG25MnCrNiMo steel. As the pressure is increased to 60 MPa, the strengthening and toughening effects of fine grains become apparent, leading to an improvement in material toughness while concurrently causing a significant decrease in wear resistance. Furthermore, the wear resistance of the material continues to improve with increasing prepared pressure, which is primarily attributed to the dominant role of solid solution strengthening on both hardness and wear resistance. Increased pressure enhances the solid solubility of alloying elements, exacerbates lattice distortion of the solid solution, generates an elastic stress field, interacts with dislocations [37], hinders the dislocation movement on slip planes, and ultimately results in increased hardness and improved wear resistance.

4. Conclusions

This study experimentally examines the effect of pressure ranging from 0 to 150 MPa on the microstructure, low-temperature toughness, hardness, and impact wear properties of ZG25MnCrNiMo steel prepared by squeeze casting. The following conclusions can be derived from this study:
(1) The samples prepared under pressure exhibited finer grains and a significant ferrite content compared to those produced without pressure. Furthermore, the secondary dendrite arm spacing of the sample produced at 150 MPa decreased by 45.3%, and the ferrite content increased by 39.1% in comparison to the unpressurized sample.
(2) The low-temperature impact toughness of the samples initially increased and then decreased as the pressure varied from 0 MPa to 150 MPa. And the toughness achieved an optimal value at a pressure of 30 MPa, which was 65.4% greater than that of gravity casting (0 MPa), while the hardness decreased by only 6.17%. The fracture was characterized by dimples, indicating a ductile fracture mechanism. With the further increase in pressure, the impact work decreased approximately linearly, and the fracture mechanism transitioned from ductile to brittle fracture.
(3) The mechanisms of pressure toughening and embrittlement in the prepared steels were as follows: at lower pressures, the significant toughening effect of fine grains and an increase in ferrite content resulted in pressure toughening; however, as pressure increased, the formation of the Widmannstatten structure and solid solution strengthening intensified the material’s embrittlement tendency, resulting in a pressure embrittlement phenomenon.
(4) The impact wear performance of the steels displayed a trend of initially decreasing and subsequently increasing across the pressure range of 0–150 MPa. The wear resistance of samples prepared without pressure and at 30 MPa was superior to that at 60 MPa, and the wear resistance deteriorated when the pressure increased to 60 MPa, after which it exhibited an upward trend as the pressure continued to rise. The wear mechanisms of the samples predominantly consisted of impact wear, adhesive wear, and minimal abrasive wear, along with notable occurrences of plastic removal, furrows, and spalling.

Author Contributions

Conceptualization, B.Q. and B.S.; methodology, B.Q.; validation, L.J., H.Y. and Z.G.; formal analysis, Z.G. and C.J.; investigation, L.J. and H.Y.; data curation, C.J. and S.L.; writing—original draft preparation, B.Q., L.J., H.Y. and S.L.; writing—review and editing, B.Q. and B.S.; project administration, B.Q.; funding acquisition, B.Q. and B.S. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Natural Science Basic Research Program of Shaanxi (2023-JC-QN-0481); the Scientific Research Program Funded by Shaanxi Provincial Education Department (23JK0489); the Science and Technology Planning Project of Bei Lin District, Xi’an (GX2316); Innovation and Entrepreneurship Training Program for College Students of Shaanxi Province (S202310702079); Science and Technology Planning Project of Xi’an (24GXFW0028) and the Young Talent Fund of Association for Science and Technology in Shaanxi, China (20240414).

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Conflicts of Interest

The authors declare no conflicts of interest.

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Figure 1. Sampling position diagram of the prepared sample.
Figure 1. Sampling position diagram of the prepared sample.
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Figure 2. Microstructure of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
Figure 2. Microstructure of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
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Figure 3. Secondary dendrite arm spacing and ferritic content of the samples prepared under different pressures.
Figure 3. Secondary dendrite arm spacing and ferritic content of the samples prepared under different pressures.
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Figure 4. SEM images at a higher magnification of the microstructure of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
Figure 4. SEM images at a higher magnification of the microstructure of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
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Figure 5. XRD analysis of the steel prepared under various pressures.
Figure 5. XRD analysis of the steel prepared under various pressures.
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Figure 6. Variation in the density and porosity of the steels prepared at different pressures.
Figure 6. Variation in the density and porosity of the steels prepared at different pressures.
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Figure 7. (a) Brinell hardness of the samples prepared under different pressures; (b) low-temperature (−40 °C) impact energy of the samples prepared under different pressures.
Figure 7. (a) Brinell hardness of the samples prepared under different pressures; (b) low-temperature (−40 °C) impact energy of the samples prepared under different pressures.
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Figure 8. Macro- and micro-impact fracture morphology of the samples prepared under different pressures: (a–a’) 0 MPa; (b–b’) 30 MPa; (c–c’) 60 MPa; (d–d’) 90 MPa; (e–e’) 120 MPa; (f–f’) 150 MPa.
Figure 8. Macro- and micro-impact fracture morphology of the samples prepared under different pressures: (a–a’) 0 MPa; (b–b’) 30 MPa; (c–c’) 60 MPa; (d–d’) 90 MPa; (e–e’) 120 MPa; (f–f’) 150 MPa.
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Figure 9. EDS analysis results of the impact fracture morphology for the sample prepared under 30 MPa: (a) site 1; (b) site 2.
Figure 9. EDS analysis results of the impact fracture morphology for the sample prepared under 30 MPa: (a) site 1; (b) site 2.
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Figure 10. (a) Relationship between wear time and wear loss of the samples prepared under different pressures; (b) wear rate of the samples prepared under different pressures.
Figure 10. (a) Relationship between wear time and wear loss of the samples prepared under different pressures; (b) wear rate of the samples prepared under different pressures.
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Figure 11. Morphology of the worn surface of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
Figure 11. Morphology of the worn surface of the samples prepared under different pressures: (a) 0 MPa; (b) 30 MPa; (c) 60 MPa; (d) 90 MPa; (e) 120 MPa; (f) 150 MPa.
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Table 1. Chemical composition of the Mn-Cr-Ni-Mo alloyed steel (wt.%).
Table 1. Chemical composition of the Mn-Cr-Ni-Mo alloyed steel (wt.%).
ElementMnCrNiMoSiCSPFe
Content1.4081.160.4450.4510.5150.2550.0200.018Balance
Table 2. Parameters of the impact wear tests.
Table 2. Parameters of the impact wear tests.
ParameterImpact EnergyImpact FrequencyRotational Speed of Counterpart Ring
Value2.5 J150 times/min150 r/min
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Qiu, B.; Jia, L.; Yang, H.; Guo, Z.; Jiang, C.; Li, S.; Sun, B. Study on the Effect of Pressure on the Microstructure, Mechanical Properties, and Impact Wear Behavior of Mn-Cr-Ni-Mo Alloyed Steel Fabricated by Squeeze Casting. Metals 2024, 14, 1054. https://doi.org/10.3390/met14091054

AMA Style

Qiu B, Jia L, Yang H, Guo Z, Jiang C, Li S, Sun B. Study on the Effect of Pressure on the Microstructure, Mechanical Properties, and Impact Wear Behavior of Mn-Cr-Ni-Mo Alloyed Steel Fabricated by Squeeze Casting. Metals. 2024; 14(9):1054. https://doi.org/10.3390/met14091054

Chicago/Turabian Style

Qiu, Bo, Longxia Jia, Heng Yang, Zhuoyu Guo, Chuyun Jiang, Shuting Li, and Biao Sun. 2024. "Study on the Effect of Pressure on the Microstructure, Mechanical Properties, and Impact Wear Behavior of Mn-Cr-Ni-Mo Alloyed Steel Fabricated by Squeeze Casting" Metals 14, no. 9: 1054. https://doi.org/10.3390/met14091054

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