Journal of Powder Metallurgy & Mining
Open Access

Titanium composites; Metallurgy of powder; Fusion of elements; Uniform microstructure; Mechanical properties

Introduction

Titanium is pursued in designing applications because of the mix of properties it gives, which incorporate great consumption obstruction, biocompatibility with the human body, high mechanical properties at high temperatures, and the most noteworthy explicit mechanical properties, as a result of the low thickness and high strength [1]. The high-temperature BCC Ti phase is stabilized at room temperature by the addition of appropriate chemical elements in + Ti alloys, which typically achieve such a combination of properties and, especially, the remarkably good balance of strength and toughness. These components, known as β stabilizers, are regularly split among isomorphous and eutectoid, contingent upon their subsequent doublestage outline with Ti. Since isomorphous elements like V, Nb, and Mo are completely soluble with Ti in both the liquid and solid states, they were the first choice for the creation of wrought Ti alloys, which require melting. The high cost of Ti alloys is exacerbated by the fact that these elements typically cost more than Ti itself. Eutectoid stabilizer elements like Fe, Mn, and Cu are typically less expensive than Ti, so they could be used to lower the alloy's overall cost [2]. In any case, they have been primarily overlooked in light of sedimentation and response issues during liquefying. Such issues can stay away from on the off chance that Ti combinations are fabricated involving powder metallurgy as they are strong state strategies [3]. In addition, powder metallurgy has other advantages, such as a high material yield and limited machining, which are significant when attempting to lower Ti's cost in comparison to that of other structural metals.

Cu and Fe are among the cheapest and most abundant eutectoid stabilizer elements. These elements can be used to make + Ti alloys and have primarily been used to make binary Ti alloys [4]. Regarding binary Ti–Cu alloys, high purity Ti and Cu powders were used to prepare Ti–xCu alloys (x = 5 and 10%; weight percentages unless otherwise noted). These alloys were ball milled for 3–6 hours, then hot pressure sintering was used to produce samples with a diameter of 40 mm under vacuum with the following parameters: 30 MPa, 120 minutes, 850–1050 °C, and furnace cooling. The samples were sintered and then extruded into 16-mm-diameter cylindrical bars at 800 °C at a rate of 10 mm/s. In another study, the same alloys were made by ball milling them for 0.5 hours, and then they were hot pressure sintered for 60 minutes in argon under vacuum conditions at 0.093 MPa and 800 °C [5]. In addition, hot pressure sintered Ti–xCu alloys with x values of 2, 5, 10, and 25% under vacuum conditions were examined, as was the possibility of producing Ti–Cu alloys through ingot casting in a vacuum non-consumable furnace, where the samples were remelted at least six times to obtain uniform compositions. A portion of the composites was likewise vacuum fixed in a gem cylinder and intensity treated at various temperatures to deliver Ti-xCu (x = 0.5, 1, 2, 5, and 10%) compounds through an argon-curve liquefying heater, where the combinations were remelted and projected into a magnesia shape at 200 °C in an outward projecting machine additionally utilized a comparative projecting strategy to get Ti-xCu (x = 2, 5, 7, and 10%) composites, which were additionally consequently heat treated at 950 °C for 3 h in a vacuum heater, trailed by cooling inside the heater to room temperature.

Laser-helped framing process has attracted expanded consideration because of its high adaptability and proficiency. Experimentation is difficult because of the interplay between temperature, microstructure, and deformation [6]. This makes it difficult to comprehend the fundamental mechanisms of the process. An efficient and preferred method for determining the intricate behavior of materials at elevated temperatures is the combination of numerical modeling and experimental observation. Thermomechanical modeling of laserassisted forming processes is the subject of numerous publications. Utilized a 3D nonlinear thermo-mechanical model, for instance, to comprehend the bending mechanism and forming characteristics of laser-assisted bending of M1A alloy, such as distortions and spring back; analyzed the strain changes during laser-assisted aluminum alloy four-point bending using FEA; taken on FE reproduction to compute the temperature and uprooting in laser-helped twisting of titanium composite and reasoned that little bowing radii could be gotten with higher laser power and additional shaping passes [7]. The above work concentrated on the misshapening conduct of materials at raised temperatures; However, the evolution of the microstructure was frequently overlooked. Thermo-metallurgical-mechanical demonstrating of laser-helped shaping cycles was seldom revealed though being read up for hot stepping and laser welding. Analyzed the residual stresses induced by phase transformation during laser beam welding of low alloy steel through a comparison of a thermo-mechanical model and a thermo-metallurgical-mechanical model declared that. For example, developed a thermo-metallurgical-mechanical model to account for the effect of boron addition and austenite deformation on transformation behavior during hot stamping The majority of studies on thermo-metallurgical-mechanical modeling were conducted to forecast post-welding or stamping residual stresses. However, despite their importance to comprehending the mechanical properties of the final formed components, extensions to the analysis of phase transformation or microstructure distribution in laser-assisted forming processes with more complex thermal passes are rarely discussed.

Method of the experiment

Mixing the required quantity of elemental powders for thirty minutes in a V-shaped blender at 45 revolutions per minute yielded Ti-xCu-xFe alloy powder blends with x values of 0.5, 1, 3.5, and 5%, respectively. After being homogenized into powder blends, a 600 MPa uniaxial pressure was used to form 40 mm specimens at room temperature [8]. The shaped Ti-xCu-xFe specimens were vacuum sintered for two hours at 1300 °C, heated to the sintering temperature at 10 °C/min, maintained at a constant vacuum level of 103 Pa, and allowed to cool inside the furnace.

Through the density before (i.e. compressibility) and after sintering, the effect of the various amounts of alloying elements on the alloy's processability was measured. In order to accomplish this, a digital caliper was used to measure the specimens' dimensions; an analytical scale was used to measure the specimens' weight; air and water were used to measure the specimens' weight; and the specific weight and density of each alloy component was used to calculate the theoretical density. The ratio of the difference between theoretical and green density values and the difference between sintered and green density values was used to calculate the densification parameter. The difference between the fully dense alloys and the sintered density was used to determine porosity.

A Kroll reactant (H2O-based solution with 2 vol. percent HF and 4 vol. percent HNO3) was used to ground, polish, and etch a sample for each of the Ti–xCu–xFe alloys in order to examine the microstructure with an Olympus BX-53 light optical microscope and a Hitachi S4700 SEM. XRD analysis (30–90° angle, 0.013° step, Cu–K source) provided additional confirmation of the phases in each Ti–xCu–xFe alloy.

The hardness of the sintered composites was estimated utilizing the Rockwell A-scale hardness (HRA) [9]. Semi static pressure strain bends of the Ti-xCu-xFe amalgams were acquired through an Instron 33-R- 4204 widespread machine. At a strain rate of 5103 s1, dogbone-shaped tensile samples with a calibrated length of 20 mm and a rectangular geometry of 2 mm x 2 mm were tested. A mechanical extensometer was used to measure the samples' deformation, and the offset method was used to calculate the yield stress. To determine the average yield stress (YS), ultimate tensile strength (UTS), and strain at fracture, at least three samples were tested for each alloy composition.

Roller forming by robots with laser assistance

MS1300 steel sheets, with a thickness of 1.0 mm, were twisted on a lab-scaled LRRF stage. The lab-scaled LRRF platform can be found in a previous publication, and the LRRF process schematic. After being clamped on the fixture, the steel sheets with a dimension of 250 mm x 60 mm were bent sequentially through three passes using laser heating and roller contact. The laser head and roller were driven simultaneously at a translational speed of 0.03 m/s by a modern robot (Kuka KR600), constrained by mathematical programming. In order to achieve preheating prior to plastic deformation, the laser beam moved 25 millimeters ahead of the roller. The laser used was a continuous wave fiber laser with a power of 1000W. Because a laser beam with the usual small spot size would have too much focused power density and damage the sheet metal surface, a rectangular laser source was used to ensure a more uniform temperature distribution at the bending corner. The spot size of the laser pillar was 4 mm × 2 mm. It should be noted that the inclination angle affects the size of the laser spot on the metal sheet, as described in Sec. 3.1. Tool steel was used to construct the 25 mm-high, 50 mm-diameter cylindrical roller.

Mechanical-thermo-metallurgical model

To account for the temperature field, microstructure evolution, and plastic deformation during the simulated LRRF process, a coupled thermo-metallurgical-mechanical model was developed and solved in Abaqus Standard. Both the stress/strain field and the microstructure field are significantly influenced by the temperature field. The thermo-metallurgical-mechanical model takes into account the major factors, including thermal expansion and temperature-induced transformation [10]. The phase transition from retained austenite to martensite is also influenced by the plastic deformation, in addition to the temperature history; In contrast, this model does not take into account the strain-induced martensitic transformation effect because martensitic steel has a negligible retained austenite fraction. The coupled thermo-metallurgical-mechanical behavior during the LRRF process is considered in the modeling procedure and framework. The overall settings, for example mathematical model, lattice, limitations, and material constitutive model, are first worked by Abaqus CAE.

To describe the laser beam's heat flux distribution and reconstruct the temperature field during LRRF, a brand-new surface heat source model is developed and integrated into Abaqus via the user subroutine DFLUX. The stage change model, representing the austenitizing conduct, austenite deterioration, and treating impact is connected to the FE model through the client subroutine USDFLD. Take note that some parameters, like the value of the phase fraction and the microhardness, are not variables in the default outputs of Abaqus; in this manner, these boundaries are named arrangement subordinate factors (SDVs) [11]. As a result, the distribution of microstructure and microhardness can be exported to the Abaqus output database file at each increment as field and history outputs. It is possible to call these SDVs in subsequent increments. The user subroutine SDVINI, which Abaqus only calls at the first solving increment, is used to set the SDVs' initial value.

MS1300 models with constitutives

An isotropic solidifying constitutive model for the MS1300 steel sheet was utilized for mechanical recreation with Abaqus. Semi static uniaxial malleable tests with the guide of a computerized picture connection (DIC) framework, allude to, were performed with a MTS widespread pliable testing machine to get the stream bend of MS1300 steel at the accompanying temperatures: 25, 200, 400, 600, and 800 degrees Celsius. At each temperature, the temperature-dependent Young's modulus and true stress vs. plastic strain curves were calculated using three different samples. The elongation is found to improve at higher temperatures, particularly when the temperature rises above 600°C. Work hardening is more common between 25°C and 200°C, while softening is more noticeable at temperatures above 400°C; The results of the tensile testing also serve as inspiration for LRRF's deformation temperature regulation.

Microstructure advancement

Since the cooling rates in LRRF are significantly higher than the critical cooling rate for martensitic transformation, the predictions made by the FE simulations that no ferrite, pearlite, or bainite will be found following LRRF processing are quite reasonable [12]. It depicts the predicted fractions of retained austenite, tempered martensite, and martensite. Take note that the fractions of retained austenite, tempered martensite, and martensite all add up to 1. Around the heating area, obvious microstructure gradients can be seen, and the deformed flange and asymmetric laser power energy applied to the clamping side still make the asymmetric microstructure distribution clear. At the outer layer of the bend, tempered martensite (95.5%) and a small amount of retained austenite (2.6%) are visible. The remaining material is martensite. The toughness of the bend is improved by the retained austenite. After forming, the inner layer of the bend's microstructure is composed of 32.2 percent tempered martensite and 67.8 percent martensite. It is also interesting to note that the middle layer has more tempered martensite (41.8 percent) than the outer and inner layers because the middle layer's tempering effect is stronger than that of the inner layer and its temperature is lower than the austenitization temperature. The tempered martensite also helps to make the bend more ductile and durable.

The bowing corner experiences the most extreme intensity input, subsequently, metallographic tests around the twisting corner were ready by cutting, mounting, crushing, cleaning, and drawing. After that, scanning electron microscope (SEM) magnifications of 2000x and 5000x were used to capture metallographic images from four distinct locations, namely the base metal and the outer, middle, and inner layers of the bend. The material that was delivered was entirely martensitic steel [13-15]. The external layer of the twist displays a refined martensitic microstructure rather than the first martensite, showing the development of new martensite. The middle layer still has a martensitic structure, and there is a lot of blocky microstructure in the grains made of ferrite and carbide precipitates (tiny white particles at a magnification of 5000x), which suggests that tempered martensite was formed. It is important to note that the steel has a low carbon content (0.21%) and is only partially tempered for a short period of time, so carbide precipitations are uncommon. Martensite and tempered martensite make up the inner layer as well. The trial microstructure circulation is steady with the mathematically anticipated results introduced.

Conclusions

The following conclusions can be drawn from this study regarding the development of new low-cost powder metallurgy ternary Ti–Cu–Fe alloys using inexpensive alloying elements like Cu and Fe:

Despite the high diffusivity of Cu and Fe, the achieved relative density values decrease with increasing addition to the Ti powder of powder particles of the alloying elements with varying morphologies and sizes. For the same sintering conditions, the amount of thermal energy used to homogenize the alloy's chemistry increases with the amount of alloying elements. However, the ternary Ti–Cu–Fe alloys' relative density values are comparable to those of other sintered Ti alloys.

In general, the addition of Cu and Fe to Ti results in the formation of a lamellar microstructure. This microstructure becomes increasingly refined in terms of prior grains and "+" lamellae as more alloying elements are added. When less than 1% of eutectoid stabilizers are added, the prior grains' morphology becomes elongated and equiaxed when more alloying elements are added. For a sufficient amount of Cu, a eutectoid substructure involving the Ti2Cu intermetallic phase within the lamellae is observed.

The sintered ternary Ti-xCu-xFe amalgams are portrayed by both flexible and plastic twisting and, for the most part, the higher how much eutectoid β stabilizers added to Ti the higher the strength and the hardness, yet the lower the capacity to endure plastic disfigurement. The yield stress/strain and ultimate tensile strength/hardness pairs are found to be in good relationships. Additionally, due to their lamellar microstructure, the sintered ternary Ti–xCu–xFe alloys' deformation and failure are controlled by the same mechanism. By the by, the degree of plastic distortion and strain solidifying is composite reliant as the science of the compound decides the stages present in the microstructure. With the gradual addition of a greater quantity of Cu and Fe, the fracture mode shifts from intergranular ductile to transgranular more brittle. The sintered ternary Ti–xCu–xFe alloys can withstand significant plastic deformation following necking at low addition levels of eutectoid stabilizers, indicating their toughness.

In this work, the LRRF cycle is created utilizing a worthwhile square shape molded laser source. A coordinated mathematical model is effectively settled through a recently evolved thermo-metallurgicalmechanical FE methodology to mimic the LRRF cycle. The temperature field, microstructure, and microhardness profile all confirm the FE simulation.

(1) The temperature field of laser-assisted forming is accurately reproduced by the newly proposed combined Gaussian-uniform surface heat source model, which takes into account the laser's inclination angle and bend profile.

(2) The coupled thermo-metallurgical-mechanical model accurately predicts the hardening effect at the outer layer caused by the formation of fresh martensite and the softening effect caused by the transition from martensite to tempered martensite.

(3) The FEM shows that the external layer of a 1.0 mm thick MS1300 steel goes through treating, austenitization, and extinguishing during LRRF while the inward layer fundamentally goes through treating. The external layer is in this manner made out of 95.5% new martensite, 2.6% held austenite, and 1.9% tempered martensite while the internal layer comprises of 67.8% starting martensite and 32.2% tempered martensite.

It is anticipated that the integrated thermo-metallurgicalmechanical finite element simulation method will also be adaptable to various manufacturing processes and additional materials (like DP and Q&P steels or titanium alloys).

Acknowledgement

None

Conflict of Interest

None

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