Effect of Heat Treatment on Tensile Properties and Microstructure of Co-Free, Low Ni-10 Mo-1.2 Ti Maraging Steel

Production of high-quality maraging steel is dependent not only on the production technology but also on the alloying design and heat treatment. In this work, cobalt-free, low nickel, molybdenum-containing maraging steel was produced by melting the raw materials in a vacuum induction melting furnace and then refining with a shielding gas electroslag remelting unit. The critical transformation temperatures of the investigated steel samples were determined experimentally by differential scanning calorimetry (DSC) analysis and theoretically aiding Thermo-Calc software. Types and chemical composition plus volume fraction and starting precipitation temperature of suggested constituents calculated with the aid of Thermo-Calc software. The microstructures of forged steel specimens that were heat-treated under several conditions were evaluated by X-ray diffraction (XRD), optical microscopy (OP), scanning electron microscopy (SEM), and electron backscattering (EBSD), in addition to transmission electron microscopy (TEM). The mechanical properties of the investigated steel specimens were evaluated by measuring the tensile strength properties and micro-hardness, furthermore, estimating their fracture surface using scanning electron microscopy at lower magnification. The metallographic results show that the microstructure of steel in aged conditions includes high-alloyed martensite and nickel-rich phase, in addition to the low-alloyed-retained-austenite, intermetallic compounds, and lavas-phase (MoCr). Furthermore, TEM and EBSD studies emphasized that the produced steel has high dislocation density with nano-sized precipitate with an average size of ~19 ± 1 nm. Moreover, the metallographic results show that the mentioned microstructure enhances the tensile properties by precipitation strengthening and the TRIP phenomenon. The tensile strength results show that the n-value of investigated steel passes two stages and is comparable with the n-value of TRIP-steel. Steel characterized by 2100 MPa ultimate tensile strength and uniform elongation of more than 7% can be produced by the investigated production routine and optimum heat treatment conditions.


Introduction
Maraging steel possesses high ultimate tensile strength and low creep under stress. This steel contains low carbon content. Hardening occurs as the martensite phase transformation arises [1] after being processed with heat treatment. Due to its high mechanical properties, this alloy has a wide application in mold casting, aerospace industries, and tool die materials. After solution treatment and the aging process in the range of 500 to 550 • C, the precipitation of intermetallic particles of fine Ni 3 (Al, Ti, Mo) and Fe 2 Mo occurs [2,3]. Strengthening occurs due to the p recipitation of the intermetallic compound particles.

Production and Material
The investigated steel was produced utilizing raw materials (steel scrap, ferroalloys, etc.). The analyzed raw materials were melted in a vacuum induction furnace. The molten steel was poured into a steel mold after completely melting and adjusting both the chemical composition and temperature. The produced ingot was used as a consumable electrode in a direct current electroslag remelting (ESR) furnace equipped with a shielding gas unit. Remelting of the consumable electrode in the electroslag remelting unit was completed under calcium fluoride-based slag 5% TiO 2 + 95% (70% CaF 2 ; 15% Al 2 O 3 ; 15% CaO). An open-air induction furnace energized at 15-kilowatt power and 450-kilohertz frequency at a temperature of approximately 1500 • C was used for preparing the master slag from CaF 2 , CaCO 3 , and Al 2 O 3 reagent-grade purchased from MERCKS. A. (Merck Limited, an affiliate of Merck KGaA, Darmstadt, Germany, 11/F, Elite Centre, 22 Hung Tong Road, Kwun Tong, Kowloon, Hong Kong, China) to produce slag similar to the chemical composition of electroslag remelting, ESR-type slag. To ensure the high homogeneity of slag powder, the dried and weighted components of the master slag components (CaF 2 , CaCO 3 , and Al 2 O 3 ) were mixed on a rotating mill for 30 min in a plastic container. Subsequently, the slag powder was collected and kept in a 250 g plastic package. Precise material balance was done for producing the investigated cobalt-free, low-nickle high-molybdenum maraging steel. Table 1 shows the chemical composition of the studied maraging steel produced by the ESR process with 10 wt.% Mo.

Thermo-Calc Studies
In this method, based on the thermodynamic calculations and knowing the chemical composition of the tested steel, both the chemical composition and the amount of each constituent were easily indicated. The multi-component system Fe-C-Ni-Mo-Cr-Ti was utilized to calculate the thermodynamic equilibria of the investigated steel composition. A sub-lattice model reported in the literature [39][40][41] described the phases in the database. Table 2 tabulates the different constituents of the investigated steel chemical composition at the end of solidification and computed in the theoretical estimations.

Differential Scanning Calorimetry (DSC)
Phase transformation temperatures of the present steel were measured utilizing DSC analysis (DSC, Labsys Setaram, Caluire, France), using temperature scans, with the rate of 0.66 • C/s in the Argon atmosphere.

Solution Treatment and Age Hardening
Based on DSC results, the solid solution treatment was carried out at 1150 • C for 2 h, followed by air cooling. Aging treatment was carried out at five different aging temperatures (400 • C, 450 • C, 500 • C, 550 • C, and 600 • C) for three aging times of 60 min, 120 min, and 240 min.

Mechanical Properties
The tensile tests were conducted according to the ASTM E08 standard at room temperature (~20 • C) until complete fracture occurred. Tensile testing apparatuses use servo hydraulic power controlled by a computer. The INSTRON 8000 ® tensile testing machine (Instron, Norwood, MA, USA) of ±100 kN load limit was used with tensile testing software, which is dedicated to the machine. An extensometer with a gauge length of 25 mm were used to measure the strain at a strain rate of 10−3/sec.
Micro-hardness analysis of the specimens was estimated along the horizontal flat plane utilizing a Vickers micro-hardness tester. During the test, a load measuring of 100 gm was employed for 20 s.

Microstructural Evaluation
Polishing of the specimen was conducted with different grades of emery paper (120, 180, 1000, and 2000), followed by cloth polishers. Then, the polished specimens were chemically etched utilizing 10% ammonium per-sulfate, then 10% nitric acid in ethyl alcohol solutions. After etching, the investigated steel sample was studied optically using a Leica optical microscope (Leica, Wetzlar, Germany) for microstructure observations.
A scanning electron microscope (SEM) Model JEOL JSM-6360 (Jeol, Tokyo, Japan) was used to analyze the surface morphology of the studied maraging steel at different experimental conditions. XRD was used to characterize the crystalline structure and phases of the tested steel. A Rigaku Ultima III X-ray diffractometer (Rigaku, Tokyo, Japan) was used to record diffraction patterns from monochromatic Cu-Kα radiation operated at 20 • C and a scan rate of 2 • /min. The X-ray machine's built-in software was used to obtain the crystallographic planes of X-ray diffraction. The volume fraction of austenite (Vγ) was specified using the following Equation (1): where Iγ is the average of the integrated intensity from the plan miller indices, such as the (111)γ and (200)γ planes, and Iα is the integrated intensity from the plan miller indices, including the (110)α planes. It was noteworthy that the value 1.4 was the correction factor in Equation (1), which was determined experimentally by many investigators [42,43]. EBSD Analysis utilized A PANalytical PANN analytical X'Pert Pro MRD goniometer (Malvern Panalytical Ltd., Malvern, UK) with a Cu tube operating at 40,000 V was employed to analyze the specimen's bulk texture. The steel micro and crystallographic texture structures were explored in the rolling (RD) and transverse direction (TD) sections. Different steps of the investigated steel specimen preparation are reported and published elsewhere [44].
A JEOL JEM 3010 TEM (Jeol, Tokyo, Japan) was utilized, which worked at a 300 kV speeding up voltage. The TEM studied circle with a thickness of 500 µm (0.5 mm) is taken from the cross portion of expelled billets (ED-plane) and precisely ground down to around 15 µm thick for TEM examinations; 3 mm plates were embedded from the specimens and cleaned to a hole utilizing a twin-stream electro-cleaning plant at 25 • C under zero degree and 15 V, using a solution of 30% nitric acid in methanol. Selected area electron diffraction (SAED) patterns were taken from a 2 µm breadth region. The specimen was cut into meager foils for TEM examination. TEM foils were cleaned at −45 • C, 10 V, in 75% (HNO 3 ) nitric acid and 25% ethanol preparation, and afterward tested in a Philips CM200 electron magnifying lens with a Gatan 300 W CCD camera.

Results and Discussion
3.1. Thermo-Calc Study Figure 1 describes the solidification behavior of the investigated steel. The solidification behavior of the studied steel specimen was calculated theoretically by the thermodynamic-equilibria model. The thermodynamic-equilibria model anticipated the chemical compositions, the volume fraction of different precipitated phases, and the starting precipitation temperature of different phases during the cooling of the studied steel. Data extracted from Figure 1 were collected and tabulated in Table 3. The data extracted from the solidification-path diagrams show that a high molybdenum content in the studied steel encourages precipitation of laves phase-C14 and retained austenite. Table 3 reports that 17.5 wt.% laves phase-C14 was precipitated at a temperature around 1090 • C, while 11 wt.% retained-austenite start precipitation at a temperature of 650 • C.
cleaned to a hole utilizing a twin-stream electro-cleaning plant at 25 °C under zero degr and 15 V, using a solution of 30% nitric acid in methanol. Selected area electron diffractio (SAED) patterns were taken from a 2 μm breadth region. The specimen was cut into me ger foils for TEM examination. TEM foils were cleaned at −45 °C, 10 V, in 75% (HNO nitric acid and 25% ethanol preparation, and afterward tested in a Philips CM200 electro magnifying lens with a Gatan 300 W CCD camera. Figure 1 describes the solidification behavior of the investigated steel. The solidific tion behavior of the studied steel specimen was calculated theoretically by the thermod namic-equilibria model. The thermodynamic-equilibria model anticipated the chemic compositions, the volume fraction of different precipitated phases, and the starting pr cipitation temperature of different phases during the cooling of the studied steel. Da extracted from Figure 1 were collected and tabulated in Table 3. The data extracted fro the solidification-path diagrams show that a high molybdenum content in the studi steel encourages precipitation of laves phase-C14 and retained austenite. Table 3 repor that 17.5 wt.% laves phase-C14 was precipitated at a temperature around 1090 °C, wh 11 wt.% retained-austenite start precipitation at a temperature of 650 °C.   The expected chemical compositions of different constituents at room temperature are presented in Figure 2. Figure 2 shows that laves phase-C14 with 56 wt.% Mo and 44 wt.% Cr were precipitated; it also shows that FCC_A2#2 carbide (titanium carbide, TiC) is the only precipitated carbide in this steel chemical composition. The dominant phase was the BCC_A2 (α-martensite), consisting predominantly of Fe and traces of other elements (Figure 2a). The second phase is the austenite (FCC_A1), consisting predominantly of nickel and a minute quantity of additional elements, as shown in Figure 2c. The laves phase (LAVES-C14) consisted mainly of Mo, Fe, Ni, Ti, and Cr. The intermetallic compounds consisted predominantly of Fe, Ti, Ni, Mo, and Si. The principal discriminator was, thus, a high Ti content, together with Si. The intermetallic compounds suggested from utilizing Thermo-Calc software were Ni 3 Ti and (Ni-Mo) 3 Ti.   (Figure 2a). The second phase is the austenite (FCC_A1), consisting predominantly of nickel and a minute quantity of additional elements, as shown in Figure 2c. The laves phase (LAVES-C14) consisted mainly of Mo, Fe, Ni, Ti, and Cr. The intermetallic compounds consisted predominantly of Fe, Ti, Ni, Mo, and Si. The principal discriminator was, thus, a high Ti content, together with Si. The intermetallic compounds suggested from utilizing Thermo-Calc software were Ni3Ti and (Ni-Mo)3Ti.

DSC Analysis
Differential scanning calorimetry, DSC diagram shows the thermic behavior of materials during heating. However, these thermic peaks are dependent on the transformation of the phases. The DSC examination of the present steel specimens was analyzed and is displayed in Figure 3. Five fundamental peaks are visible on the DSC diagram of the studied steel specimen. However, the first exothermic peak typically happened due to the martensite recovery and the retained-austenite evolution [39][40][41]. The second exothermic peak appears with continued heating of the investigated steel, which accelerates the decomposing of the intermetallic compound, such as Ni3(Ti, Mo) [45]. Ni3(Ti, Mo) are representative in clusters from 400 to 550 °C, according to Tewari et al. [45].
The third exothermic peak is typically visible due to phase transformation from martensite to austenite. This transformation occurs at a temperature range from 600 to 700 °C, depending on the molybdenum contents in the studied steel. Furthermore, the last two peaks are compared to the interaction between the structure and the circumstances of the solid solution. In the high temperature, peaks appear due to the dissociation of the precipitate laves phase-C14. The DSC results were interpreted with the aid of Thermo-Calc studies in the prior section, Figure 1 and Table 3. The DSC curve for the studied steel specimen shows a few shifts at high temperatures, which could be induced by the effect of impurities, such as oxides or sulfides.
The maraging steel was aged (tempered) at a temperature range of 400-600 °C, and the solution heat treatment temperature was > 1083 °C, according to the DSC curve. It is noteworthy that the studied steel contains high alloying element contents; thus, microsegregation dominated. Therefore, the studied steel requires high homogenized temperatures. Differential scanning calorimetry results recommended that the solution heat

DSC Analysis
Differential scanning calorimetry, DSC diagram shows the thermic behavior of materials during heating. However, these thermic peaks are dependent on the transformation of the phases. The DSC examination of the present steel specimens was analyzed and is displayed in Figure 3. Five fundamental peaks are visible on the DSC diagram of the studied steel specimen. However, the first exothermic peak typically happened due to the martensite recovery and the retained-austenite evolution [39][40][41]. The second exothermic peak appears with continued heating of the investigated steel, which accelerates the decomposing of the intermetallic compound, such as Ni 3 (Ti, Mo) [45]. Ni 3 (Ti, Mo) are representative in clusters from 400 to 550 • C, according to Tewari et al. [45].
The third exothermic peak is typically visible due to phase transformation from martensite to austenite. This transformation occurs at a temperature range from 600 to 700 • C, depending on the molybdenum contents in the studied steel. Furthermore, the last two peaks are compared to the interaction between the structure and the circumstances of the solid solution. In the high temperature, peaks appear due to the dissociation of the precipitate laves phase-C14. The DSC results were interpreted with the aid of Thermo-Calc studies in the prior section, Figure 1 and Table 3. The DSC curve for the studied steel specimen shows a few shifts at high temperatures, which could be induced by the effect of impurities, such as oxides or sulfides.
The maraging steel was aged (tempered) at a temperature range of 400-600 • C, and the solution heat treatment temperature was >1083 • C, according to the DSC curve. It is noteworthy that the studied steel contains high alloying element contents; thus, microsegregation dominated. Therefore, the studied steel requires high homogenized temperatures. Differential scanning calorimetry results recommended that the solution heat treatment temperature be increased to more than 1083 • C, in this work, to ensure a more successful heat treatment. treatment temperature be increased to more than 1083°C, in this work, to ensure a m successful heat treatment.
The results extracted from the DSC diagrams show that the high molybdenum tent in studied steel has a remarkable effect in different transformation temperature shown in Figure 3. The previous results were confirmed by aided Thermo-Calc studie in the prior section and Figure 1.

Aging Treatment
Thermo-calc and DSC studies recommended that the studied steel requires high mogenization temperatures so that the solution heat treatment temperature be incre to 1150 °C in this work to ensure a more successful heat treatment. The results extracted from the DSC diagrams show that the high molybdenum content in studied steel has a remarkable effect in different transformation temperatures, as shown in Figure 3. The previous results were confirmed by aided Thermo-Calc studies, as in the prior section and Figure 1.

Aging Treatment
Thermo-calc and DSC studies recommended that the studied steel requires high homogenization temperatures so that the solution heat treatment temperature be increased to 1150 • C in this work to ensure a more successful heat treatment. Figure 4 shows the hardness variation of the investigated steel with aging time and temperature. The Figure 4 result is explained by an increase in the hardness that may be attributed to the intermetallic compounds strengthening mechanism of the steel sample during aging. The hardness profile of the studied steel samples increased gradually with increasing aging temperature up to the peak hardness, as shown in Figure 4. The peak hardness appeared at 500 • C. Increasing temperature beyond the peak hardness temperature led to a sharp decrease in the hardness values. This decrease in hardness occurs due to the dissociation of the intermetallic and laves-phase, as explained in the Thermo-Calc studies section. Zhang et al. [46] stated that as the aging temperature rises, the volume fraction of blocky retained-austenite in martensite blocks increases. In addition, the retained-austenite phase deteriorates the studied steel toughness. The peak hardness value indicates the optimum requirement of the aging process (temperature and time). From Figure 4, the authors decided the optimum parameters of the heat treatment process to get the designed mechanical properties. Table 4 shows the optimum heat treatment conditions for the production of the studied steel; it is anticipated that tensile and elongation properties will be enhanced as the material ages at optimum conditions. tained-austenite phase deteriorates the studied steel toughness. The peak hardness v indicates the optimum requirement of the aging process (temperature and time). F Figure 4, the authors decided the optimum parameters of the heat treatment proce get the designed mechanical properties. Table 4 shows the optimum heat treatment ditions for the production of the studied steel; it is anticipated that tensile and elonga properties will be enhanced as the material ages at optimum conditions.  The solid-solution-annealing (austenitization) temperature of 1150 °C was assu to be sufficient to prevent the formation of any precipitates at some stage in homogen tion annealing or soaking during forging operations. Therefore, all of the lavas-p would precipitate during the subsequent cooling after finishing the heat treatment ( solution annealing) or forged-processing. Conversely, the precipitation of the rest o intermetallic compounds, such as Ni3Ti, Ni3(Ti-Mo), Ni3Mo, etc., should only occur du the aging heat treatment for all three conditions. The thermo-calc studies recommen the mentioned heat treatment temperatures. The forging and homogenization, in add to the aging temperatures, are summarized in Table 4. Furthermore, the forging and mogenization, in addition to the aging cycle, are schematically represented in Figure The solid-solution-annealing (austenitization) temperature of 1150 • C was assumed to be sufficient to prevent the formation of any precipitates at some stage in homogenization annealing or soaking during forging operations. Therefore, all of the lavas-phase would precipitate during the subsequent cooling after finishing the heat treatment (solid solution annealing) or forged-processing. Conversely, the precipitation of the rest of the intermetallic compounds, such as Ni 3 Ti, Ni 3 (Ti-Mo), Ni 3 Mo, etc., should only occur during the aging heat treatment for all three conditions. The thermo-calc studies recommended the mentioned heat treatment temperatures. The forging and homogenization, in addition to the aging temperatures, are summarized in Table 4. Furthermore, the forging and homogenization, in addition to the aging cycle, are schematically represented in Figure 5.   Figure 6 shows the optical micrographs of the present steel after ST and in the peakaged conditions at 500 °C for different times. The microstructure in the peak aged conditions consisted of martensite packets within prior-austenite grains. The transformed-austenite-grains were distinguished because of the preferential etching along their boundaries. Therefore, the martensite-packets within the austenite grain did not extend beyond the respective prior-austenite grain boundary.  Figure 6 shows the optical micrographs of the present steel after ST and in the peakaged conditions at 500 • C for different times. The microstructure in the peak aged conditions consisted of martensite packets within prior-austenite grains. The transformed-austenitegrains were distinguished because of the preferential etching along their boundaries. Therefore, the martensite-packets within the austenite grain did not extend beyond the respective prior-austenite grain boundary.  The substructure of the martensite could not be monitored because of the fineness of the martensite laths. During the aging of the steels under investigation, the well-known precipitation reactions lead to hardening. The initial precipitation in cobalt-free molybdenum-containing maraging steel at 480 °C occurs as Ni3Mo, which on prolonged aging is replaced by either lava-phase (Fe2Mo) or the σ phase (FeMo) [4]. Since the alloy additionally contains titanium as an auxiliary hardener, the precipitation of Ni3Ti has also been reported; alternatively, it has been suggested that part of the titanium may be present in the molybdenum precipitate, i.e., as Ni3(Mo, Ti). This microstructure suggested that the mechanical properties of the studied steel were enhanced due to the well-known substructure of lathe martensite consisting predominantly of a high density of dislocation tangles within the laths [13,14]. Figure 7 shows XRD patterns for the studied heat-treated maraging steel samples. Furthermore, Figure 7 emphasizes the present microstructure consisting of martensite, retained austenite, and different precipitated particles. Figure 7 shows that the highest intensity peak was present in the patterns corresponding to the martensite (body-centered cubic) phase for all the studied steel samples. Figure 7 shows that the produced maraging steel primarily consists of the martensite phase. The martensite formed because of the phase transformation process and alloying elements' contents. The X-ray diffraction pattern of the austenite peaks in the studied heat-treated steels reflects the high stability of the retained austenite, as shown in Figure 7. The high stability of the retained austenite The substructure of the martensite could not be monitored because of the fineness of the martensite laths. During the aging of the steels under investigation, the wellknown precipitation reactions lead to hardening. The initial precipitation in cobalt-free molybdenum-containing maraging steel at 480 • C occurs as Ni 3 Mo, which on prolonged aging is replaced by either lava-phase (Fe 2 Mo) or the σ phase (FeMo) [4]. Since the alloy additionally contains titanium as an auxiliary hardener, the precipitation of Ni 3 Ti has also been reported; alternatively, it has been suggested that part of the titanium may be present in the molybdenum precipitate, i.e., as Ni 3 (Mo, Ti). This microstructure suggested that the mechanical properties of the studied steel were enhanced due to the well-known substructure of lathe martensite consisting predominantly of a high density of dislocation tangles within the laths [13,14]. Figure 7 shows XRD patterns for the studied heat-treated maraging steel samples. Furthermore, Figure 7 emphasizes the present microstructure consisting of martensite, retained austenite, and different precipitated particles. Figure 7 shows that the highest intensity peak was present in the patterns corresponding to the martensite (body-centered cubic) phase for all the studied steel samples. Figure 7 shows that the produced maraging steel primarily consists of the martensite phase. The martensite formed because of the phase transformation process and alloying elements' contents. The X-ray diffraction pattern of the austenite peaks in the studied heat-treated steels reflects the high stability of the retained austenite, as shown in Figure 7. The high stability of the retained austenite could be attributed to the highly alloyed element contents. Additional small peaks correspond to the nickel intermetallic compounds and laves-phase, such as  Figure 7. The exact type of the precipitate particles could not be identified from the XRD patterns since the primary peaks were close together. However, all phase peaks with small intensities in a 2-theta angle ranging between 20 • and 40 • were expected to represent nickel intermetallic compounds or laves-phase.  Figure 7. The exact type of the precipitate parti not be identified from the XRD patterns since the primary peaks were close However, all phase peaks with small intensities in a 2-theta angle ranging be and 40° were expected to represent nickel intermetallic compounds or laves-ph However, these small phase peaks were significantly noticed in the aged 240 min, but the same peaks with lower intensities were also noticeable in all studied steel samples. No additional peak occurred that could be related to t precipitates, probably due to their small volume fraction.

XRD Analysis
The amount of retained austenite formed after solution treatment and ag optimum condition for the present steel was calculated utilizing X-ray diffracti tion (1)) and tabulated in Table 5. Table 5 shows that the quantity of retained au the studied steel gradually increased with increasing aging time. The quantity o austenite increased due to the coarsening of the intermetallic compounds, which the alloying element from the martensite phase. Furthermore, diffusion of th elements from the martensite phase produces martensite with lower alloying that decrease the stability of the martensite phase and lead to its transformati tenite. This can explain the decrease of hardness with increasing aging tempe shown in Figure 4.  However, these small phase peaks were significantly noticed in the aged sample at 240 min, but the same peaks with lower intensities were also noticeable in all the other studied steel samples. No additional peak occurred that could be related to the Ti-rich precipitates, probably due to their small volume fraction.
The amount of retained austenite formed after solution treatment and aging at the optimum condition for the present steel was calculated utilizing X-ray diffraction (Equation (1)) and tabulated in Table 5. Table 5 shows that the quantity of retained austenite of the studied steel gradually increased with increasing aging time. The quantity of retained austenite increased due to the coarsening of the intermetallic compounds, which consume the alloying element from the martensite phase. Furthermore, diffusion of the alloying elements from the martensite phase produces martensite with lower alloying elements that decrease the stability of the martensite phase and lead to its transformation to austenite. This can explain the decrease of hardness with increasing aging temperature, as shown in Figure 4.

EBSD Analysis
The microstructural morphologies observed in the investigated steel corresponded well with the lath-martensite (blue area) structure characterized by EBSD analyses. Figure 8 shows the EBSD IQ-phase map of the experimental sample. The larger blocky reversed γ (red area) formed in martensite blocks are displayed in Figure 8. Moreover, the amount of reversed-austenite was measured to be 13.2 vol.%, while the laves-phase C14 (green area) was measured to be 17 vol.%. The obtained results confirmed the X-ray diffraction test. The black lines indicate the high-angle grain boundary (HAGB) over 15 • , which mainly includes the hierarchical boundaries prior-austenite grains, packets, and blocks [47], Figure 8b. Meanwhile, the misorientation below 5 • reveals sub-grains colored in red. The microstructural morphologies observed in the investigated steel corresponded well with the lath-martensite (blue area) structure characterized by EBSD analyses. Figure  8 shows the EBSD IQ-phase map of the experimental sample. The larger blocky reversed γ (red area) formed in martensite blocks are displayed in Figure 8. Moreover, the amount of reversed-austenite was measured to be 13.2 vol.%, while the laves-phase C14 (green area) was measured to be 17 vol.%. The obtained results confirmed the X-ray diffraction test. The black lines indicate the high-angle grain boundary (HAGB) over 15°, which mainly includes the hierarchical boundaries prior-austenite grains, packets, and blocks [47], Figure 8b. Meanwhile, the misorientation below 5° reveals sub-grains colored in red.  Figure 9 presents the EBSD result of the M10AC steel specimen aged at 500 °C for 120 min, revealing the fine austenite (γ) grains were distributed in a ferrite (α-martensite) matrix. The retained γ phases often appeared along the grain boundaries in the α-martensite matrix. Many of the fine-γ grains were oriented along the Z-direction (Figure 9b), forming a {001} texture of the retained γ phase (Figure 9).  Figure 9 presents the EBSD result of the M10AC steel specimen aged at 500 • C for 120 min, revealing the fine austenite (γ) grains were distributed in a ferrite (α-martensite) matrix. The retained γ phases often appeared along the grain boundaries in the αmartensite matrix. Many of the fine-γ grains were oriented along the Z-direction (Figure 9b), forming a {001} texture of the retained γ phase (Figure 9). The corresponding stereographic projections of Figure 10a depicting the low-index orientations obtained by performing the EBSD analyses are displayed in Figure 10a-d. EBSD analysis revealed the orientation relationship between the fine retained γ grains oriented along the Z-direction and the adjacent α grains in the martensite structure (Figure 10a). As exhibited in the obtained stereographic projections (Figure 10a,b), the fine γ  The corresponding stereographic projections of Figure 10a depicting the low-index orientations obtained by performing the EBSD analyses are displayed in Figure 10a-d. EBSD analysis revealed the orientation relationship between the fine retained γ grains oriented along the Z-direction and the adjacent α grains in the martensite structure ( Figure 10a). As exhibited in the obtained stereographic projections (Figure 10a,b), the fine γ grain had an orientation relation of (ı ı ı)γ//(0 ı ı)α and [ī 0 ı]γ//[īī ı]α concerning the adjacent α grain. The stereographic projections (Figure 10c,d) representing the fine γ grain also had a different variant of (ı ı ı)γ//(0 ı ı)α and [ī 0 ı]γ//[īī ı]α, concerning another adjacent α grain. The determined orientation relationship corresponds to the Kurdjumov-Sachs (K-S) orientation relationship between lath martensite and austenite and is consistent with the crystallographic features of the lath martensite structure in conventionally quenched maraging steels. (c) (d) Figure 10. Stereographic projection of the 111γ /011α and 011γ /111α poles. Figure 11 denotes the histogram plot of 10 wt.% Mo AC aged-steel specimen. Furthermore, Figure 11 shows small secondary precipitates of Mo introduced in the final microstructure, with an average diameter of 19 ± 1 nm and narrower size distribution.
In summary, the experimented steel specimen showed that the Mo-rich phase precipitates were present, with distinctively different sizes. The laves-phase precipitate (MoCr) was observed in the microstructure and found within the martensite laths and along the lath boundaries.  Figure 11 denotes the histogram plot of 10 wt.% Mo AC aged-steel specimen. Furthermore, Figure 11 shows small secondary precipitates of Mo introduced in the final microstructure, with an average diameter of 19 ± 1 nm and narrower size distribution.
In summary, the experimented steel specimen showed that the Mo-rich phase precipitates were present, with distinctively different sizes. The laves-phase precipitate (MoCr) was observed in the microstructure and found within the martensite laths and along the lath boundaries.

TEM Analysis
From the TEM-image, given in Figures 12 and 13, a high dislocation density was realized in the α-martensite matrix in both solid solution and aged steel specimen. It is clear that the dislocation density decreased while the martensite lath boundaries gradually disappeared with the film-like reversed γ observed at the inter-laths. The above results indicate that the austenite reversion occurred with increased aging time. The precipitate size on average, was found to be approximately 40 nm in length and 2.5 nm in thickness. The intermetallic compound precipitates, though entirely small, were found to show a variety of fringe contrast that could be seen clearly in the dark field images (Figure 12b) obtained by using precipitate reflections. For specific electron beam directions, these fringes were found to be parallel to the length of the intermetallic compounds precipitates. In other intermetallic compounds, however, the fringes were found to be perpendicular to their intermetallic compound's length.

TEM Analysis
From the TEM-image, given in Figures 12 and 13, a high dislocation density was realized in the α-martensite matrix in both solid solution and aged steel specimen. It is clear that the dislocation density decreased while the martensite lath boundaries gradually disappeared with the film-like reversed γ observed at the inter-laths. The above results indicate that the austenite reversion occurred with increased aging time. The precipitate size, on average, was found to be approximately 40 nm in length and 2.5 nm in thickness. The intermetallic compound precipitates, though entirely small, were found to show a variety of fringe contrast that could be seen clearly in the dark field images (Figure 12b) obtained by using precipitate reflections. For specific electron beam directions, these fringes were found to be parallel to the length of the intermetallic compounds precipitates. In other intermetallic compounds, however, the fringes were found to be perpendicular to their intermetallic compound's length.   Figure 14 represents the SAED patterns of the intermetallic compound precipitate in this steel sample, which could be indexed in codes of the hexagonal, eta (η) phase-type; η Ni3Ti having lattice parameters a = 0.5101 nm and c = 0.8307 nm. Since this phase has been seen to contain Mo, it will be designated as the Ni3(Ti, Mo) phase. In addition to the rodshaped precipitates, the investigated steel sample showed the presence of nearly spherical shape-precipitates, as shown in Figure (12b). Selected area diffraction patterns (SADP) from these precipitated particles could be indexed in the codes of the hexagonal Fe2Mo phase (Figure 14).  Figure 14 represents the SAED patterns of the intermetallic compound precipitate in this steel sample, which could be indexed in codes of the hexagonal, eta (η) phase-type; η Ni 3 Ti having lattice parameters a = 0.5101 nm and c = 0.8307 nm. Since this phase has been seen to contain Mo, it will be designated as the Ni 3 (Ti, Mo) phase. In addition to the rod-shaped precipitates, the investigated steel sample showed the presence of nearly spherical shape-precipitates, as shown in Figure 12b. Selected area diffraction patterns (SADP) from these precipitated particles could be indexed in the codes of the hexagonal Fe 2 Mo phase ( Figure 14).  Figure 14 represents the SAED patterns of the intermetallic compound precipitate in this steel sample, which could be indexed in codes of the hexagonal, eta (η) phase-type; η Ni3Ti having lattice parameters a = 0.5101 nm and c = 0.8307 nm. Since this phase has been seen to contain Mo, it will be designated as the Ni3(Ti, Mo) phase. In addition to the rodshaped precipitates, the investigated steel sample showed the presence of nearly spherical shape-precipitates, as shown in Figure (12b). Selected area diffraction patterns (SADP) from these precipitated particles could be indexed in the codes of the hexagonal Fe2Mo phase (Figure 14).

Effect of Aging Time on Tensile Properties
Figures 15 and 16 illustrate the engineering and true stress-strain curves obtained for Mo-containing maraging steel (10 wt.% Mo), respectively. Figure 15a represents the engineering stress-strain curves for solid solution steel specimens, while, Figure 15b-d represents engineering stress-strain curves for aged steel specimens at 500 °C for 60 min, 120 min, and 240 min, respectively. Figure 15 shows that the steel specimen aged at any aging time shows a remarkable increase in all tensile properties except the ductility (elongation, ef %). The mentioned result, which may be attributed to during aging of the intermetallic compounds, show that coarsening and accumulation led to precipitation of the hardening strengthening process. This strengthening mechanism is diffusion-controlled, depending mainly on the temperature and time. It was noteworthy that the effect of aging time on tensile properties is a concern in this manuscript. Furthermore, the microstructure instability leads to an increase in the emergence rate of several constituents, such as low-alloyed martensite and high-alloyed retained austenite, in addition to the low-alloyed austenite phase. The decrease in the elongation % (ef %) of the tested steel may be due to the high prompt formation, precipitation, and coarsening of intermetallic compounds. Moreover, austenite and martensite in maraging steel have high dislocation density. Precipitation of intermetallic compounds diminishes the dislocation movement, which causes the strengthening of the investigated steel through the particles-dislocations interaction mechanism. The blocking of dislocation movement decreased the ductility (ef %) of the studied steel. On the other hand, non-metallic inclusions may be considered as sites of crack formation and initiation. Therefore, the presence of non-metallic inclusions diminishes the tensile properties of the investigated steel, especially the elongation %, because it has a role as an obstacle of dislocation movement and site of crack initiation and propagation.
To confirm the previous phenomena, the tensile fractography of the solid solution and the aged M10AC sample (500 °C/120 min) at room temperature is presented in Figure  17. The dimple mode of the fracture on the fracture surface of tensile tested specimens revealed a ductile failure mode. Dimples were prominent in all samples tested. From the fractographic analysis, we concluded that ductile fracture was dominant, and many micro-void formations took place.

Effect of Aging Time on Tensile Properties
Figures 15 and 16 illustrate the engineering and true stress-strain curves obtained for Mo-containing maraging steel (10 wt.% Mo), respectively. Figure 15a represents the engineering stress-strain curves for solid solution steel specimens, while, Figure 15b-d represents engineering stress-strain curves for aged steel specimens at 500 • C for 60 min, 120 min, and 240 min, respectively. Figure 15 shows that the steel specimen aged at any aging time shows a remarkable increase in all tensile properties except the ductility (elongation, e f %). The mentioned result, which may be attributed to during aging of the intermetallic compounds, show that coarsening and accumulation led to precipitation of the hardening strengthening process. This strengthening mechanism is diffusion-controlled, depending mainly on the temperature and time. It was noteworthy that the effect of aging time on tensile properties is a concern in this manuscript. Furthermore, the microstructure instability leads to an increase in the emergence rate of several constituents, such as low-alloyed martensite and high-alloyed retained austenite, in addition to the low-alloyed austenite phase. The decrease in the elongation % (e f %) of the tested steel may be due to the high prompt formation, precipitation, and coarsening of intermetallic compounds. Moreover, austenite and martensite in maraging steel have high dislocation density. Precipitation of intermetallic compounds diminishes the dislocation movement, which causes the strengthening of the investigated steel through the particles-dislocations interaction mechanism. The blocking of dislocation movement decreased the ductility (e f %) of the studied steel. On the other hand, non-metallic inclusions may be considered as sites of crack formation and initiation. Therefore, the presence of non-metallic inclusions diminishes the tensile properties of the investigated steel, especially the elongation %, because it has a role as an obstacle of dislocation movement and site of crack initiation and propagation.   To confirm the previous phenomena, the tensile fractography of the solid solution and the aged M10AC sample (500 • C/120 min) at room temperature is presented in Figure 17. The dimple mode of the fracture on the fracture surface of tensile tested specimens revealed a ductile failure mode. Dimples were prominent in all samples tested. From the fractographic analysis, we concluded that ductile fracture was dominant, and many micro-void formations took place. It was noteworthy that the changes in both the microstructure and chemical composition of different constituents led to an increase in the transformation potential of the austenitic phase to the martensitic phase. Strengthening may have occurred due to the TRIP (transformation induced plasticity) phenomenon during deformation in the tensile test [46,47]. This assumption was confirmed utilizing the computation of the strain hardening exponent (n-value) for the tested samples. Figure 18 reported that all investigated steel showed two stages according to the strain %. The first stage (at low strain %) was for strain accumulation, while the second stage was for phase transformation (induced martensite); additionally, beyond the second stage, instability or necking began. All aged specimens showed a higher strain hardening exponent than solid solution steel specimens at all strain %. The aging process slightly increased the uniform-elongation from 4.31 to 5.47% (>25%) compared to the solid solution process. Due to the precipitation of a high quantity of intermetallic compounds and retained austenite phase. On the other hand, total elongation % (ef %) was reduced from 11.4 to 7.64%. Furthermore, the aging process significantly improved the ultimate tensile strength for all investigated aging times, as shown in Table 6. It was noteworthy that the changes in both the microstructure and chemical composition of different constituents led to an increase in the transformation potential of the austenitic phase to the martensitic phase. Strengthening may have occurred due to the TRIP (transformation induced plasticity) phenomenon during deformation in the tensile test [46,47]. This assumption was confirmed utilizing the computation of the strain hardening exponent (n-value) for the tested samples. Figure 18 reported that all investigated steel showed two stages according to the strain %. The first stage (at low strain %) was for strain accumulation, while the second stage was for phase transformation (induced martensite); additionally, beyond the second stage, instability or necking began. All aged specimens showed a higher strain hardening exponent than solid solution steel specimens at all strain %. The aging process slightly increased the uniform-elongation from 4.31 to 5.47% (>25%) compared to the solid solution process. Due to the precipitation of a high quantity of intermetallic compounds and retained austenite phase. On the other hand, total elongation % (ef %) was reduced from 11.4 to 7.64%. Furthermore, the aging process significantly improved the ultimate tensile strength for all investigated aging times, as shown in Table 6.
Investigated steels offer additional advantages derived from a unique multiphase microstructure and add retained austenite; in addition, they added intermetallic compounds to the martensite microstructure. The tensile deformation encourages the retained-austenite phase transformation into martensite. The transformation process enhances the strength of the investigated steel, which was explained the effect of the TRIP phenomenon [48,49]. Therefore, this transformation continues with additional deformation as long as sufficient retained austenite is present, allowing the investigated steel to maintain a high average n-value of 0.09235 to 0.1914 over a tensile test, as shown in Figure 18.   Investigated steels offer additional advantages derived from a unique multiphase microstructure and add retained austenite; in addition, they added intermetallic compounds to the martensite microstructure. The tensile deformation encourages the retained-austenite phase transformation into martensite. The transformation process enhances the strength of the investigated steel, which was explained the effect of the TRIP phenomenon [48,49]. Therefore, this transformation continues with additional deformation as long as sufficient retained austenite is present, allowing the investigated steel to maintain a high average n-value of 0.09235 to 0.1914 over a tensile test, as shown in Figure 18.
The authors suggest that the mentioned characteristic allows for forming more complex geometries and potentially at a reduced thickness to perform the mass reduction. Furthermore, during the deformation of the steel specimen, the retained austenite remaining in the microstructure can subsequently transform into martensite during the event of The authors suggest that the mentioned characteristic allows for forming more complex geometries and potentially at a reduced thickness to perform the mass reduction. Furthermore, during the deformation of the steel specimen, the retained austenite remaining in the microstructure can subsequently transform into martensite during the event of the collision. This result makes the investigated steels an appropriate choice for parts in crush zones on a vehicle.
Meanwhile, the n-value for solid solution steel specimen was~0.09235, the n-value for aged steel specimens ranged 0.1601-0.1914, which is comparable to values for TRIP steel (n-value = 0.23) with the same ultimate tensile strength value. It is well-known that the n-value of the material, primarily steel, depends intensely on the different chemistries and technological processing paths used by steelmakers. The n-value is a function of the strain accumulation of the deformation that affects the transformation of the retained austenite. It is well-known that a higher n-value reduces strain gradients, allowing for more complex deformation. Lower n-value concentrates strains, leading to early failure.
Since different locations in deformed steel pursue several strain tracks with varying amounts of deformation, the n-value for the designed steel could alter with both parts design and location. The modified microstructures of the advanced high-strength steel permit different property relationships to enhance each steel type and grade to particular application requirements. Figure 19 shows the work hardening rate-true strain diagram. The work hardening rates of the investigated steel at both solid solution and aged condition dropped quickly after yielding, followed by decreasing continuously at higher strains. The solid-solution and aged samples showed similar behavior of work hardening. Figure 19 shows that aged steel samples were characterized by a higher work hardening rate, which is thought to be related to the large fraction of deformed structure. It is speculated that the additional work hardening is affected by the mechanical incompatibility in the plastic strain gradient and heterogeneous structure evolution multiaxial stresses [48,49] during the elastic-plastic transition period. The previous results showed that the investigated steel strengthened due to the austenite transforming into martensite during the deformation process (tensile test). Moreover, the strengthening of investigated steel was explained by the transformationinduced plasticity (TRIP) phenomenon.
application requirements. Figure 19 shows the work hardening rate-true strain diagram. The work hardening rates of the investigated steel at both solid solution and aged condition dropped quickly after yielding, followed by decreasing continuously at higher strains. The solid-solution and aged samples showed similar behavior of work hardening. Figure 19 shows that aged steel samples were characterized by a higher work hardening rate, which is thought to be related to the large fraction of deformed structure. It is speculated that the additional work hardening is affected by the mechanical incompatibility in the plastic strain gradient and heterogeneous structure evolution multiaxial stresses [48,49] during the elastic-plastic transition period. The previous results showed that the investigated steel strengthened due to the austenite transforming into martensite during the deformation process (tensile test). Moreover, the strengthening of investigated steel was explained by the transformation-induced plasticity (TRIP) phenomenon. The true stress-true strain curves are drawn as in Figure 16 for verifying the instability point, according to the Consideré criterion [50] for necking. It was established that the solid solution steel specimen had low uniform strain (4.31%) and slowly fractured after necking. On the other hand, a considerable elongation (from a true strain of 5.06% to 5.47%) was realized in the aged steel specimens by postponing early localized necking and/or resisting cracking. For more clarity to differentiate between the tested cases, Figure 20 is constructed.
It is noteworthy that the resultant data was limited to Co-free low-nickel maraging steel produced by the double-melting technique vacuum melting followed by electroslag remelting. In addition, the resulting data is limited for this material after forging and heat-treated at the optimum condition demonstrated in the main text of the manuscript. Therefore, to complete the mechanical properties investigation, studying the behavior of the investigated steel at high temperatures needs more attention. Moreover, suggesting the appropriateness of the investigated steel alloy for working as a creep resistance material is the central subject of our further investigations.
The true stress-true strain curves are drawn as in Figure 16 for verifying the instability point, according to the Consideré criterion [50] for necking. It was established that the solid solution steel specimen had low uniform strain (4.31%) and slowly fractured after necking. On the other hand, a considerable elongation (from a true strain of 5.06% to 5.47%) was realized in the aged steel specimens by postponing early localized necking and/or resisting cracking. For more clarity to differentiate between the tested cases, Figure  20 is constructed. It is noteworthy that the resultant data was limited to Co-free low-nickel maraging steel produced by the double-melting technique vacuum melting followed by electroslag remelting. In addition, the resulting data is limited for this material after forging and heattreated at the optimum condition demonstrated in the main text of the manuscript. Therefore, to complete the mechanical properties investigation, studying the behavior of the investigated steel at high temperatures needs more attention. Moreover, suggesting the appropriateness of the investigated steel alloy for working as a creep resistance material is the central subject of our further investigations.

Conclusions
The effect of heat treatment (i.e., solid solution annealing followed by air cooling in addition to aging at variable temperature and time) on tensile properties and microstructure of air-cooled Co-free, low Ni-10 Mo-1.2 Ti maraging steel was investigated. From the presented results, the following conclusions may be summarized.
1. The low carbon content of the investigated steel leads to very low or no carbide precipitate. High molybdenum contents encouraged lava-phase precipitate with chemical composition formula MoCr and enforced the investigated steel to precipitate η(Ni3Ti). 2. The microstructure of aged steel specimens consists of (i) high-alloyed martensite phase, (ii) nickel-rich phase, (iii) low-alloyed retained-austenite phase, (iv) intermetallic compounds, (v) and lavas-phase (MoCr). The presence of this microstructure enhances the tensile properties by precipitation strengthening and TRIP phenomenon. 3. TEM and EBSD studies emphasized that the produced steel has high dislocation density with nano-sized precipitate with an average size ~19 ± 1 nm.

Conclusions
The effect of heat treatment (i.e., solid solution annealing followed by air cooling in addition to aging at variable temperature and time) on tensile properties and microstructure of air-cooled Co-free, low Ni-10 Mo-1.2 Ti maraging steel was investigated. From the presented results, the following conclusions may be summarized.
1. The low carbon content of the investigated steel leads to very low or no carbide precipitate. High molybdenum contents encouraged lava-phase precipitate with chemical composition formula MoCr and enforced the investigated steel to precipitate η(Ni 3 Ti). 2. The microstructure of aged steel specimens consists of (i) high-alloyed martensite phase, (ii) nickel-rich phase, (iii) low-alloyed retained-austenite phase, (iv) intermetallic compounds, (v) and lavas-phase (MoCr). The presence of this microstructure enhances the tensile properties by precipitation strengthening and TRIP phenomenon. 3. TEM and EBSD studies emphasized that the produced steel has high dislocation density with nano-sized precipitate with an average size~19 ± 1 nm. 4. EBSD studies revealed that the fine austenite (γ) grains were distributed in a ferrite (αmartensite) matrix. The retained γ phases often appeared along the grain boundaries in the α-martensite matrix. Many of the fine γ grains were oriented along the Z-direction, forming a {001} texture of the retained γ phase. 5. Calculation of the strain hardening exponent (n-value) shows that investigated steel (passing two stages) and its n-value are comparable with the n-value of TRIP steel. Aging the steel specimens at 500 • C and 120 min shows a high work-hardening rate compared with other aged steel specimens.