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Article

Formation of Heterogeneous Nucleation of B2-NiAl in Hot Rolled Fe-Mn-Al-C Plate: A Novel Composition and Processing Route for Lightweight High Strength Steel Containing Nickel

by
Michael Piston
1,*,
Laura Bartlett
1,
Krista R. Limmer
2,
Daniel M. Field
2 and
Billy C. Hornbuckle
2
1
Department of Metallurgical Engineering, Missouri University of Science and Technology, Rolla, MO 65401, USA
2
DEVCOM Army Research Laboratory, APG, Adelphi, MD 21005, USA
*
Author to whom correspondence should be addressed.
Metals 2024, 14(12), 1342; https://doi.org/10.3390/met14121342
Submission received: 10 September 2024 / Revised: 4 November 2024 / Accepted: 21 November 2024 / Published: 26 November 2024
(This article belongs to the Special Issue Development of Advanced High-Strength Steels)
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Abstract

In this study, a novel lightweight Fe-Mn-Al-C steel composition and thermomechanical processing route was developed to produce a fully austenitic microstructure with a uniform intragranular dispersion of B2-NiAl precipitation in order to overcome the significant challenge of strengthening hot-rolled Fe-Mn-Al-C steels while retaining toughness. The new composition and processing methods allow for the processing of ultrahigh-strength Fe-Mn-Al-C steel containing nickel as thicker gauge plate for a multitude of new automotive and structural applications where lightweighting is critical. The composition investigated in this study was a fully austenitic Fe-21Mn-9Al-1C-8Ni wt% steel. Two hot rolling methods were investigated: the first procedure involved lower temperature rolling cycles to precipitate B2-NiAl during hot rolling and reheating. The second method involved higher temperature rolling to precipitate B2-NiAl after thermomechanical processing during a short isothermal treatment. The lower temperature rolling produced plate with an ultimate tensile strength of 1120 MPa and a Charpy V-Notch (CVN) toughness of 24 J at −40 °C. After the high temperature rolling procedure, precipitation of B2-NiAl through a subsequent precipitation hardening step resulted in reduced B2-NiAl size and improved the ultimate tensile strength above 1300 MPa. The two novel processing routes of a single composition can be performed with current manufacturing capabilities to produce hot rolled plate strengthened by B2-NiAl precipitation to various hardness (ranging from 33 to 41 HRC) and strength levels (ranging from 1100 to 1320 MPa ultimate tensile strength) while retaining 22–27% elongation.

1. Introduction

Fe-Mn-Al-C steels continue to be of great interest in many applications because of a high strength to weight ratio compared to conventional steels. The significant density reduction is the result of aluminum additions of up to 12 wt% by reducing the density of the austenite and ferrite phases through direct substitution of Fe with Al and by dilation of the austenite lattice [1,2,3]. Carbon additions up to 2 wt% in the Fe-Mn-Al-C system have been shown to be four times more effective at reducing density as Al [1,4]. The combined effect results in densities up to 17% lower than traditional quenched and tempered martensitic steels.
Depending on alloying additions, Fe-Mn-Al-C steels can range from fully ferritic to fully austenitic. Fully austenitic Fe-Mn-Al-C steels typically yield the best mechanical property combinations and are the target microstructure for this study. These fully austenitic Fe-Mn-Al-C steels generally have ultimate tensile strengths (UTS) and total elongations ranging from 800 to 1500 MPa and 30 to 80%, respectively [5,6,7,8,9,10,11,12,13,14,15,16,17,18,19,20]. Fully ferritic Fe-Mn-Al-C steels have reported UTS values between 200 and 600 MPa, and total elongations of 10–40% [5,21,22,23,24,25,26,27]. Impact toughness in these systems is infrequently reported, but it is of great interest and benefit for applications that require high energy absorption.
High-manganese austenitic steels exhibit a variety of active deformation mechanisms depending on the stacking fault energy (SFE). Stress-induced γ → ε transformation known as transformation induced plasticity (TRIP) occurs when SFE is below 20 mJ/m2, twinning induced plasticity (TWIP) occurs by forming deformation twins at SFE between 20 and 40 mJ/m2, and partial to perfect dislocation glide will occur as the SFE becomes greater than 40 mJ/m2 [28,29]. The deformation of <20 mJ/m2 SFE face-centered cubic (FCC) metals is typically dominated by planar glide with widely spaced partial dislocations, difficulty of cross slip, and increased friction stress [30,31,32]. High manganese steels with slightly higher SFE in the range of 20–40 mJ/m2 also display this effect, forming dislocation cells and planar dislocation structures at true strains below 0.1. However, these steels will begin to exhibit TWIP at true strains of 0.3–0.4 [30]. SFE in Fe-Mn-Al-C steels has been shown to be strongly increased through additions of Al [5,33,34]. The 9 wt% Al steel in the current study has a calculated SFE of 80.2 mJ/m2 without Ni using the thermodynamic model of Yang et al. [35]. The addition of Ni is predicted to further increase the SFE according to first principles calculations by Limmer et al. [36]. However, the presence of short-range ordering in the solid solution austenite has been reported as the driving factor for planar glide, regardless of the higher SFE [6]. Therefore, the high SFE in combination with short-range ordered zones will result in a deformation mechanism dominated by planar glide in high aluminum Fe-Mn-Al-C steels.
Steels containing more than 0.7 wt% C and 7 wt% Al can be strengthened by the precipitation of coherent nanosized κ-carbide, (Fe,Mn)3AlC, during aging within typical precipitation hardening temperature ranges of 450 °C to 600 °C [37]. The strengthening contribution of κ-carbide is reported to be between 100 and 300 MPa for volume fractions between 0.05 and 0.35 [38]. Unfortunately, Fe-Mn-Al-C steels that are precipitation hardened in the peak aged condition exhibit significant losses in work hardening rate and toughness. This is due to dislocation shearing of κ-carbides creating an easy pathway for dislocation glide along the slip plane, known as the glide plane softening effect [5,7,8,9]. Previous research has shown that this localized deformation can produce cracking along slip planes and a phenomenon called glide plane decohesion that produces a cleavage-like fracture with characteristic slip traces on fracture surfaces [39]. Therefore, strengthening through κ-carbide alone reduces the ability of aged Fe-Mn-Al-C steels to achieve high strength and good combinations of both ductility and toughness.
A study by Kim et al. has shown that another method to increase strength and hardness in Fe-Mn-Al-C steels results from controlled additions of Ni [40]. These nickel additions have been shown to produce a hard intermetallic phase of ordered B2-NiAl in duplex Fe-15Mn-10Al-0.8C-5Ni steel. The following morphologies of B2-NiAl were observed after cold rolling and subsequent precipitation hardening: Type 1 was δ-ferrite rolled into stringers, Type 2 was intergranular B2-NiAl, and Type 3 was intragranular B2-NiAl [40]. After a two-minute heat treatment at 900 °C, a fine dispersion of 50–300 nm intragranular B2-NiAl was heterogeneously nucleated on shear bands within the deformed austenite. Ultimate tensile strengths up to 1600 MPa and total elongations of 20% were reported from these Ni-strengthened steels [40]. It was also reported that the B2-NiAl act as Orowan barriers that may limit the glide plane softening effect in high SFE steels [40]. An alternative method to precipitate B2-NiAl was reported by Zargaran et al. [41], where pre-existing κ-carbide precipitates greater than 10 nm can also act as nuclei for intragranular B2-NiAl.
A separate study was conducted by Piston et al. on a Fe-18Mn-10Al-0.9C-5Ni steel to explore the feasibility of plate product [42]. It was found that the composition and processing method, which excluded cold rolling, produced 200–1000 nm B2-NiAl precipitates at grain boundaries and primary δ-ferrite stringers leading to brittle fracture at room temperature. In duplex Fe-Mn-Al-C steels with Ni, δ-ferrite stringers and grain boundary B2-NiAl have been shown to be deleterious to Charpy V-notch (CVN) toughness, promoting preferential crack paths when fracture occurs parallel to the δ-ferrite stringers. In order to reduce this detrimental effect, the current study was performed to develop a new composition of fully austenitic Fe-Mn-Al-C-Ni steels and processing methodology that address the issues of low CVN toughness and precipitation of intragranular B2-NiAl without cold rolling.

2. Materials and Methods

A 72.6-kg heat with the nominal composition of Fe-21Mn-9Al-1C-8Ni (in wt%) was prepared in a coreless induction furnace and melted under argon gas cover. The melt was then calcium treated with 98% Ca iron-cored wire for a total of 23.8 g of Ca and tapped into a teapot ladle at 1560 °C. The resulting steel melt was then poured into a bottom-gated phenolic resin-bonded silica sand Y-block mold. The runner contained a 10 ppi ceramic filter to minimize oxide bifilms in the casting. Chemistry results were acquired by optical emission spectroscopy (OES) using an OXFORD FOUNDRYMASTER. Carbon and sulfur concentrations were measured using gas combustion analysis with a LECO CS600 C/S analyzer. The final dimensions extracted from the casting were 50 mm in width × 82.5 mm in length × 63.5 mm in thickness obtained from the base of the Y-blocks immediately above a steel chill plate. The ingot was heated to 1200 °C and held at temperature for 2 h before rolling to dissolve the B2-NiAl, κ-carbide, and δ-ferrite present in the as-cast structure and to reduce segregation.
Thermodynamic modeling was performed in ThermoCalc 2024b using the TCFE14 database. This was done to attempt to identify a single phase austenite region for solution treatment and the B2-NiAl stability temperature. Two hot rolling procedures were performed on a rolling mill with a two-high configuration to investigate the effect of thermomechanical processing on B2-NiAl shape and size that started by solutionizing the ingot at 1200 °C for 2 h. Both procedures used hot rolling reductions of 1.02 mm per pass at 12 RPM with 127-mm diameter rolls. A rolling cycle will be defined as four passes totaling 5.08 mm followed by a 5 min reheat to maintain the target temperature range. In procedure 1 (P1), rolling cycles occurred between 1050 and 900 °C to a final thickness of 12.7 mm. Upon reaching the target gauge thickness, the plate was immediately reheated to 1050 °C for three minutes, followed by water quenching. In procedure 2 (P2), rolling cycles occurred between 1200 and 1050 °C to a thickness of 17.8 mm (72% hot reduction). The plate was transferred to a furnace at 1050 °C for five minutes followed by a rolling cycle between 1050 and 900 °C to a final thickness of 12.7 mm followed by agitated water quenching. The P2 plate was subsequently precipitation hardened at 950 °C for 10 min. Schematics for the P1 and P2 conditions are provided in Figure 1.
Specimens for optical metallography were sectioned with an abrasive saw and polished down to a 0.05 µm finish using standard polishing procedures. Etching was performed with a 10% nital solution. Scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDS) were performed on a Helios NanoLab 600 at an accelerating voltage of 30 kV and working distance of 4.5 mm. A ThermoFisher Scientific Helios G4 UX DualBeam microscope was used for thinning at 30 kV with a final step at 5 kV, and transmission electron microscopy (TEM) was performed on an JEOL 2100F operated at 200 kV. ImageJ V1.54g was used to measure precipitate size on a minimum of three regions. A minimum of three sub-sized tensile specimens for each processing route with 6 ± 0.1 mm width, 5.95 ± 0.05 mm thickness, and 25 ± 0.1 mm gauge length were prepared parallel to the rolling direction by Electrical Discharge Machining (EDM) and tested at room temperature per ASTM E8 with a strain rate of 0.01 s−1. Charpy V-notch toughness was determined with a minimum of three specimens per orientation using 10 mm × 10 mm × 55 mm (full size) specimens at −40 °C in the L-T and T-L orientations in accordance with ASTM E23.

3. Results

The chemical analysis of the Fe-21Mn-9Al-1C-8Ni (wt%) steel is presented in Table 1. Thermodynamic equilibrium calculations were utilized from the composition in Table 1 to predict phase transformations to guide heat treatments and rolling temperatures. As shown in Figure 2, equilibrium calculations predict formation of δ-ferrite as the primary phase during solidification, with the liquid combining with the δ-ferrite to form austenite. The calculated solidus for the steel was 1240 °C and B2-NiAl was predicted to be stable up to the solidus. The solution treatment temperature of 1200 °C was selected to dissolve the highest fraction of B2-NiAl from the casting as possible without producing incipient melting. At 1200 °C, thermodynamic modeling predicts 32.6 wt% B2-NiAl, which corresponds to 32.9 vol% B2-NiAl. However, after solution treatment of the as-cast steel at 1200 °C, the microstructure in Figure 3 contains 1.53 ± 0.35 area% of B2-NiAl within interdendritic regions and on grain boundaries. While not directly comparable, it is a strong indicator that the thermodynamic modeling simulation over-predicts the amount of both δ-ferrite and B2-NiAl present in this composition.
After the P1 hot rolling between 1050 and 900 °C, the specimen in Figure 4a,b contains a microstructure consistent with non-recrystallized elongated austenite grains decorated with intergranular B2-NiAl and intragranular B2-NiAl particles within all austenite grains. Segregation banding evidenced as light and dark etching regions are observed after the 2 h at 1200 °C solution treatment and the subsequent thermomechanical processing. The secondary electron image is presented in Figure 4b of the P1 processed steel, and shows at higher resolution that the B2-NiAl have an elongated spheroid morphology within the matrix and an average length of 559 ± 240 nm and width of 252 ± 100 nm with an aspect ratio of 2.2, as reported in Table 2.
The P2 rolling procedure performed the majority of the reduction between 1200 and 1050 °C to induce recrystallization followed by one set of strain building passes between 1050 and 900 °C. The hot-rolled and water-quenched microstructure is shown in Figure 5a. The steel in this condition shows evidence of recrystallization that was not observed in the lower temperature rolling procedure, and intragranular B2-NiAl precipitation was also not present. The microstructure of the P2 processed plate after precipitation hardening at 950 °C for 10 min is presented in Figure 5b,c. It is characterized by a fine dispersion of B2-NiAl plates within the austenite matrix, as well as the presence of B2-NiAl on grain boundaries. The intragranular B2-NiAl precipitate size, as shown in Table 2, had an average length of 124 ± 56 nm and thickness of 7.8 ± 3.5 nm, with an aspect ratio of 15.9, which is seven times greater than the P1 processed steel. A 0.69 ± 0.13 µm precipitate-free zone was observed along austenite grain boundaries. The corresponding scanning transmission electron microscope (STEM) bright field (BF) images of the P2 specimen are provided in Figure 6a,b. Selected area diffraction patterns (SADP) confirm the matrix is austenite with an average lattice parameter of 3.615 ± 0.016 Å, and the precipitates are ordered B2-NiAl with an average lattice parameter of 2.884 ± 0.028 Å. No κ-carbide was observed in the P2 condition.
In order to understand the influence of processing on mechanical properties, the stress–strain response is shown in Figure 7a,b, with the calculated work hardening rate as a function of true strain in Figure 7c. For comparison, the investigated steel is shown relative to the previously reported hot rolled duplex 5% Ni steel by the author showing that austenitic hot rolled steels can achieve similar or improved strength, ductility, and work hardening compared to hot rolled duplex steels [42]. The mechanical properties are summarized in Table 3, in addition to B2-NiAl morphology. The P1 steel has the lowest strength and total elongation, while the P2 steel has a yield strength of 940 MPa, the highest UTS of 1320 MPa, and the greatest total elongation. It is interesting to note that the P2 steel also exhibits the highest work-hardening rates as shown in Figure 7c. The strain hardening exponent was determined by the Ludwik–Hollomon equation [43], and all steels exhibited non-linear behavior. The P2 steel ranged from n = 0.10 at the early stages of plastic deformation to a linear n = 0.27 in the 10–20% plastic deformation range, as shown in Figure 7d. Both the P1 processed steel and the reference 5 wt% Ni steel had a linear n = 0.30 within this range.
The tensile performance of the P2 condition had a 21% increase in yield strength to 940 MPa, and a 20% increase in UTS to 1320 MPa, compared to P1 without any loss of total elongation. The P1 specimen did not contain δ-ferrite/B2-NiAl stringers, which resulted in a significantly improved T-L to L-T ratio of 0.64 when compared to prior work by the author on the 5 wt% Ni steel that had a toughness ratio of 0.26 and similar hardness. In the L-T orientation, the P1 processed steel displayed excellent toughness of greater than 37 J at −40 °C. The P2 condition had higher strength, however, there was a 68% reduction in the L-T notch toughness of 12 J compared to P1. Observations of the Charpy fracture surface of P2 in Figure 8a revealed brittle fracture occurring in planes of a grain, as indicated by white arrows. The tensile fracture surface in the P2 condition in Figure 8b also contained this brittle fracture, even under quasi-static strain rates.

4. Discussion

In the current study, a novel Fe-21.4Mn-8.9Al-1.03C-7.7Ni composition was used to produce a lightweight, austenitic, ultrahigh-strength, and hot rolled steel with up to 1320 MPa in strength and ductility greater than 26%. The previous Fe-18.3Mn-10.1Al-0.9C-4.5Ni steel (5% Ni) studied by the author had problems with primary δ-ferrite [42]. The alloy was redesigned with additional austenite stabilizers such as manganese and carbon, and additional nickel was predicted to significantly reduce the amount of primary δ-ferrite for the 8% Ni steel and increase the stability of B2-NiAl ordering in the ferrite closer to solidification temperatures in Figure 9. Thermodynamic equilibrium calculations predicted the solidus for the 5 and 8 wt% Ni steel to be 1265 °C and 1240 °C, respectively. The solution treatment temperature was selected to be 40 °C below the solidus to dissolve the most B2-NiAl from the casting as possible without melting the interdendritic regions. At 1200 °C, thermodynamic modeling predicts 46 wt% disordered ferrite in the 5% Ni steel and 32 wt% B2-NiAl in the 8% Ni steel. This is a contradiction of observations of the as-cast, solution-treated at 1200 °C microstructures in Figure 3 which contain 1.53 ± 0.35 percent surface area of primary δ-ferrite or B2-NiAl in the 8% Ni steel, which is significantly lower than expected, based on equilibrium calculations. A limitation of the thermodynamic models is a significant overprediction of the δ-ferrite/B2-NiAl equilibrium stability in the 8% Ni steel. From Figure 3, a 2-h solution treatment at 1200 °C dissolved almost all B2-NiAl, which is a significant deviation from the predicted equilibrium amount of 32 wt% B2-NiAl, as shown in Figure 9. Kim et al. [44] has also pointed out that experimental information for phase equilibria in quaternary systems such as Fe-Mn-Al-C is lacking, and thermodynamic models do not always accurately predict phase fractions in these complex and high-alloy steel systems. Decomposition of the austenite into B2-NiAl and κ-carbide is predicted to occur at 850 °C in the 5 wt% Ni steel. The higher Mn and lower Al increased the stability of austenite and reduced the predicted decomposition temperature below 770 °C in the 8% Ni steel. Therefore, hot rolling and isothermal heat treatments should be maintained above these temperatures to avoid coarse κ-carbide precipitation on grain boundaries.
The B2-NiAl precipitates in the P1 specimen were formed during the hot rolling and reheat cycles between 900 and 1050 °C, as shown in Figure 4a,b. The microstructure is consistent with non-recrystallized austenite with a uniform dispersion of intragranular B2-NiAl precipitates formed by the hot rolling of specimen P1 which had an average length of 559 ± 240 nm and a width of 252 ± 100 nm. The large standard deviation in measured precipitate size was due to B2-NiAl forming at varying roll and reheat cycles. The final three-minute hold at 1050 °C in P1 is theorized to be responsible for forming the finest observed intragranular B2-NiAl precipitation in Figure 4b as this newly nucleated B2-NiAl precipitation would have undergone the fewest rolling and reheat cycles. In contrast, the largest B2-NiAl precipitates were a result of nucleation in early rolling and reheat cycles, and are expected to have coarsened. A limited amount of B2-NiAl was observed on γ-γ grain boundaries. An increase in B2-NiAl precipitation density within the banded regions was observed in the former interdendritic regions of the P1 specimen in Figure 4a. It is important to note that precipitation of B2-NiAl through this method occurs by rolling and heat treating exclusively between 900 and 1050 °C below the recrystallization temperature without requiring cold rolling or isothermal treatments after water quenching. Therefore, reducing the total number of roll and reheat cycles in future work will greatly decrease the B2-NiAl precipitate size.
In contrast, hot rolling procedure 2 (P2) was designed to hot roll the majority of the reduction without forming B2-NiAl between 1200 and 1050 °C. The final 5.08 mm of reduction was performed from 1050 to 900 °C to introduce the necessary accumulated strain prior to water quenching for B2-NiAl to nucleate upon. The as-rolled microstructure of the P2 specimen consisted of an austenitic matrix with recrystallized grains and intergranular B2-NiAl present as shown in Figure 5a. This unintended precipitation on grain boundaries was likely a result of the 1050 °C reheat for 5 min prior to the strain building passes where insufficient dislocations were present to nucleate intragranular B2-NiAl. Deformed annealing twins were also observed from deformation induced from the final passes between 1050 and 900 °C. Future investigations would benefit from controlled rolling between 900 and 1200 °C without a reheat cycle to prevent an isothermal hold in a temperature range capable of precipitating B2-NiAl without heterogeneous nucleation sites present from accumulated strain.
After precipitation hardening for 10 min at 950 °C, the intragranular B2-NiAl precipitation within the austenite grains of specimen P2 in Figure 5b was too fine to resolve optically. The morphology of the B2-NiAl was confirmed as plates growing in preferential orientations within a grain as shown in Figure 5c. TEM diffraction patterns presented in Figure 6a,b confirmed that nano-sized precipitates of plate/disc morphology are ordered B2-NiAl with an average lattice parameter of 2.884 ± 0.028 Å. This is consistent with the work of Jiang et al. who reported a lattice parameter of 2.882 ± 0.012 Å for B2-Ni(Al,Fe) in a maraging steel [45]. The secondary electron images in Figure 5c were used to determine that the P2 condition produced B2-NiAl precipitates with an average length of 124 ± 56 nm and a thickness of 7.8 ± 3.5 nm. Kim et al. reported a comparable size for intragranular B2-NiAl size of 50–300 nm after cold rolling to 1 mm and precipitation hardening at 900 °C for 2 min [40]. Therefore, cold rolling is not a requirement to nucleate intragranular B2-NiAl precipitation. However, the short precipitation hardening steps used in cold rolled sheet to achieve peak strength are not applicable to thicker plate section sizes. The rapid growth of the intragranular B2-NiAl was observed to occur along slip bands due to these bands acting as significant pipe diffusion pathways [46]. The activation energy for diffusion of nickel along these pathways has been reported to be 0.4–0.7 times the activation energy for Ni diffusion through bulk FCC [46]. Reducing the growth kinetics and size of B2-NiAl during precipitation hardening of hot rolled plate can be simply achieved by reducing the treatment temperature below 950 °C while staying above the κ-carbide solvus to avoid nucleating carbide on grain boundaries.
In the current study, the entire rolling procedure for all steels occurred above 900 °C followed by agitated water quenching to avoid the formation of κ-carbide as confirmed by TEM observations in Figure 6a,b. As the κ-carbide solvus was calculated to be 850 °C for the 5% Ni steel and 770 °C for the 8% Ni steel, the B2-NiAl is not expected to have formed on pre-existing κ-carbide nuclei greater than 10 nm in size, as observed in the work of Zargaran et al. [41]. Although there is limited literature available for this system, most authors report fine intragranular B2-NiAl precipitates on shear bands in non-recrystallized austenite only after cold rolling and subsequent annealing in the temperature range of 900 to 1000 °C [40]. This work shows that a uniform dispersion of B2-NiAl can form when rolling and reheats occur between 900 and 1050 °C in P1 where κ-carbide is not predicted to be stable. Additionally, higher temperature rolling can induce the recrystallization of austenite grains before strain is accumulated at lower temperatures.
The available literature on mechanical properties of nickel-containing Fe-Mn-Al-C steels are also limited, and compositions are typically in the range of Fe-(15–21Mn)-10Al-(0.8–1)C-5Ni. Steels within this range are reported as containing δ-ferrite, and increasing the precipitation hardening times and temperatures rapidly coarsens B2-NiAl, resulting in a significant reduction in strength [40,41,42,47]. Previous studies have shown that increasing the precipitation hardening temperature from 870 to 900 °C during 15 min of soaking resulted in a decrease in the UTS from 1470 MPa to 1350 MPa in a Fe-15Mn-10Al-0.8C-5Ni steel [40]. Therefore, for applications requiring greater plate thickness, the potential issue arises of non-uniform B2-NiAl coarsening at the surface during typical convective heating. Kim et al. additionally showed that increasing the precipitation hardening time at 900 °C from 2 to 15 min reduced the UTS from 1550 to 1350 MPa and increased the total elongation from 20.3 to 31.8% in cold rolled 1-mm thick sheet [40]. In our present study, the P2 steel that was precipitation hardened for 10 min at 950 °C has a comparable UTS of 1320 MPa with a similar total elongation of 26.6%, without requiring cold work to nucleate B2-NiAl, as well as avoiding the detrimental presence of δ-ferrite on T-L orientation toughness as a fully austenitic composition [42].
The tensile curves from each condition in the current study displayed continuous yielding and strain hardening behavior consistent with what has been observed in Fe-12Mn-8Al-0.8C duplex steel, as well as a similar composition of Fe-16Mn-9.6Al-4.9Ni-0.86C steel [48,49]. In addition, the work hardening rate as a function of true strain in the previously mentioned studies displayed similar behavior and three distinct stages, as observed for the current steels in Figure 7b. Initially, the strain hardening rate rapidly decreased after the onset of plastic deformation (stage A), as shown in Figure 7c. In stage B, the strain hardening rate decreased slightly or even plateaus before decreasing abruptly again before fracture in stage C [48]. Wu et al. compared the strain hardening behavior of austenitic, duplex, and triplex Fe-Mn-Al-C steels, and showed that the initial strain hardening rates (at ε = 5%) of duplex steels (γ+α) and triplex steels (γ+α+κ) are much higher, ranging from 1350 MPa for an austenitic Fe-28Mn-10Al-1.0C steel to 2250 MPa for a triplex Fe-12Mn-8Al-0.8C with 21% ferrite [48]. The high initial work hardening rate was attributed to the concentration of deformation and dislocations at the interface between hard ferrite and soft austenite, similar to what is observed in dual-phase steels [48,50]. The 5 and 8 wt% Ni-containing steels have much higher initial work hardening rates (at ε = 5%) of between 2900 and 3600 MPa, respectively, than what has been observed for austenitic and duplex Fe-Mn-Al-C steels without Ni addition. The P2, P1, and reference 5 wt% Ni steel had a strain hardening exponent (n) of 0.27, 0.30, and 0.30, respectively. The P2 condition with n = 0.27 at 940 MPa yield strength was in good agreement with the work done by Kim [49] who reported a strain hardening exponent for a duplex Fe-16Mn-10Al-0.86C-5Ni of n = 0.26 at ~1000 MPa yield strength. The B2-NiAl-containing lightweight steels displayed a decreasing strain hardening exponent as yield strength increased due to the activation of less favorable slip systems at higher stress [49].
Yang et al. reported a detailed study of the strain hardening behavior of a Fe-16Mn-10Al-0.86C-5Ni steel using in situ synchrotron diffraction to measure the lattice strains in austenite and B2-NiAl [51]. They attributed the extraordinarily high strain hardening rate to the high back stresses that developed from incompatibility of stress–strain partitioning between the austenite and B2-NiAl [51]. However, subsequent studies by Kim [49] of a Fe-16Mn-10Al-0.86C-5Ni austenitic steel with and without B2-NiAl showed similar strain hardening behavior. Their study suggested that the austenite matrix chemistry and short-range ordering in the highly alloyed FCC lattice plays a larger role in strain hardening than the hard dislocation barrier B2-NiAl as plate/discs. The comparable strain hardening exponents between the austenitic Fe-21Mn-9Al-1C-8Ni steel in this study and the duplex Fe-16Mn-10Al-0.86C-5Ni steel from Kim [49] agrees with their findings and strongly indicates that the matrix plays the dominant role in strain hardening mechanism for these systems, regardless of the presence of B2-NiAl stringers ordered from δ-ferrite or intragranular B2-NiAl precipitation. This is likely due to planar dislocation motion on the {111} austenite planes and ease of cross slip to new {111} planes with minimal resistance to dislocation motion passing an incoherent precipitate like B2-NiAl on the same plane.
The specific yield strength (SYS) as a function of ductility of the currently studied steels (solid diamonds) in comparison to hot and cold rolled and precipitation hardened Ni-containing B2-NiAl-strengthened steels of similar compositions (solid circles) and other austenitic and duplex Fe-Mn-Al-C steels are presented in Figure 10 [40,42,51,52,53,54,55,56,57,58]. It is shown that nickel-containing steels strengthened by B2-NiAl have comparable or better combinations of strength and ductility than Fe-Mn-Al-C alloys without Ni, and can achieve greater ultimate tensile strengths while retaining toughness. Hot rolled nickel-free Fe-Mn-Al-C steels without B2-NiAl typically achieve tensile strengths of up to 1200 MPa, as reported by Acselrad et al. for a wrought Fe-29Mn-8.8Al-1C-1.33Si steel in the peak aged condition [59]. Although strength greatly increases, age hardening via the precipitation of κ-carbide has been shown to be extremely detrimental to ductility and notch toughness [60,61]. Bartlett et al. report an increase in strength from 960 to 1160 MPa, but a precipitous loss of elongation from 56 to 14% in a Fe-30Mn-9Al-1Si-0.9C-0.5Mo steel after aging for 60 h at 530 °C [60]. Therefore, the strength and ductility of the B2-NiAl-strengthened 8% Ni steel in the current study outperforms austenitic Fe-Mn-Al-C steels that are strengthened by age hardening and κ-carbide precipitation without requiring lengthy aging treatments, but does not achieve the increased strength levels from B2-NiAl-strengthened Fe-Mn-Al-C processed as cold rolled sheet with extremely short precipitation hardening times [37,51].
Impact toughness is a significantly under-reported mechanical property in the nickel-containing Fe-Mn-Al-C system. To the author’s knowledge, Charpy V-notch properties have not been reported thus far for austenitic Ni-containing B2-NiAl-strengthened Fe-Mn-Al-C steels. In the current study, anisotropy in notch toughness was greatly improved by eliminating primary δ-ferrite. The P1 condition was fully austenitic with a UTS of 1100 MPa with 37 J in the L-T and 23.7 J in the T-L orientation at −40 °C, as shown in Table 3. In a prior study by the author, a duplex Fe-18Mn-10Al-0.9C-5Ni steel had a comparable UTS of 1120 MPa with 28.5 J in the L-T and 7.3 J in the T-L orientation at −40 °C [48]. This significant anisotropy was attributed to crack propagation along the length of δ-ferrite stringers ordered to B2-NiAl in the T-L orientation. Unlike most austenitic steels, Fe-Mn-Al-C steels have been reported to have a ductile to brittle transition at low temperatures caused by brittle crack nucleation at the intersection of planar slip-induced dislocation defects on {111} and tensile separation of glide planes producing cleavage-like facets [62,63].
The steel in the present study was calculated to have an SFE above 80 mJ/m2. Higher SFE metals like this are reported to contain narrow partial dislocations that can cross-slip the screw component onto another plane, resulting in a dislocation structure known as wavy glide that improves ductility but reduces capacity to work harden [30,64]. However, the presence of planar glide in these Fe-Mn-Al-C high SFE steels is attributed to a preferential dislocation motion through sheared weak dislocation barriers like short range ordered zones and/or coherent and shearable precipitate such as κ-carbide [7,14,31,49,65]. The result of this dislocation motion in Figure 8 is apparent in both the (a) quasi-static tensile strain rate and (b) the high strain rate Charpy impact test fracture surface. Planar fracture of individual grains was observed in both the impact test and tensile test fracture face, though significantly more prevalent in the higher strain rate Charpy impact test. When compared to austenitic Fe-Mn-Al-C steels that are strengthened by κ-carbide, the 8% Ni steel exhibits improved notch toughness at equivalent strength levels. As shown in Table 3, the P1 specimen achieved a hardness of 32.6 HRC and a UTS of 1100 MPa with 37 J in the L-T orientation at −40 °C. This can be compared with a UTS of 1050 MPa and a notch toughness of 5 J at room temperature (20 °C), reported by Hale et al. for a Fe-30Mn-8Al-1C steel strengthened by κ-carbide precipitation [66]. While no tensile properties were acquired, Ley et al. [67] reported a hot rolled Fe-30Mn-9Al-1C-1Si-0.7Mo steel strengthened by κ-carbide that achieved a hardness level of 350 HBW (~35.5 HRC) with a T-L impact toughness at room temperature (22 °C) of 15 J. Therefore, nickel-containing fully austenitic Fe-Mn-Al-C steels strengthened with B2-NiAl exhibited improved toughness over duplex nickel-containing steels, or those strengthened by κ-carbide.

5. Conclusions

An austenitic Fe-21Mn-9Al-1C-8Ni steel composition and two thermomechanical processing routes were developed to produce intragranular B2-NiAl without cold rolling that can open up avenues for production of plate from lightweight steels with nickel. It was determined that the use of CALPHAD to predict phase stability in these systems is still limited. It was shown that it is possible to precipitate B2-NiAl uniformly within austenite grains in 12.7-mm thick plate by hot rolling above 900 °C. This has only previously been achieved through either cold rolling or using a pre-existing nuclei such as κ-carbide in sheet steel less than 3 mm thick. The hot rolled material in this study showed comparable strength and elongation to similar duplex cold rolled steel for a given precipitation time and temperature, but conventional heat treatment of thicker plate sections are limited in their maximum achievable strength due to the rapid coarsening of the plate B2-NiAl precipitates. Decreasing B2-NiAl size significantly increased yield strength but played a smaller role in strength contribution from strain hardening than expected, indicating the matrix plays the most important role in strain hardening for these materials. Impact properties in the P2condition were decreased due to the higher volume fraction of intragranular B2-NiAl, the presence of undesirable intergranular B2-NiAl, and individual austenite grains exhibited cleavage fracture along planes that was more prevalent at higher strain rates. Future improvements to the system would be to further decrease the isothermal treatment temperature below 950 °C but above the κ-carbide stability temperature, slightly reduce the amount of nickel in the system to avoid interdendritic B2-NiAl after solution treatment, and increase solution treatment time to further reduce banding. The processing methodology can be improved with larger reductions per pass and fewer reheats. For hot rolling between 1050 and 900 °C (P1), a larger reduction per pass produced finer and more uniform distributions of B2-NiAl precipitates. In the steel rolled between 1200 and 1050 °C followed by a reheat at 1050 °C and four passes between 1050 and 900 °C (P2), larger reductions per pass could bypass the need for a 1050 °C reheat and avoid intergranular B2-NiAl, which is deleterious to toughness.

Author Contributions

Conceptualization, M.P.; methodology, M.P.; software, M.P.; validation, M.P. and L.B.; formal analysis, M.P.; investigation, M.P. and B.C.H.; resources, M.P. and L.B.; data curation, M.P.; writing—original draft preparation, M.P.; writing—review and editing, M.P., L.B., K.R.L., D.M.F. and B.C.H.; visualization, M.P., L.B., K.R.L. and D.M.F.; supervision, L.B.; project administration, L.B. and K.R.L.; funding acquisition, L.B. All authors have read and agreed to the published version of the manuscript.

Funding

Research was sponsored by the DEVCOM ARL Army Research Office and was accomplished under Co-operative Agreement Number W911NF-20-2-0251. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of the DEVCOM ARL Army Research Office or the U.S. Government. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation herein.

Data Availability Statement

The datasets presented in this article are not readily available due to security clearance requirements.

Conflicts of Interest

The authors declare no conflicts of interest.

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Figure 1. Diagram of thermomechanical processing and subsequent heat treatment for the 8% Ni steel for (a) procedure 1 (P1) and (b) procedure 2 (P2). Temperature regions observed in this study for strain accumulation and recrystallization (RXN) are labelled.
Figure 1. Diagram of thermomechanical processing and subsequent heat treatment for the 8% Ni steel for (a) procedure 1 (P1) and (b) procedure 2 (P2). Temperature regions observed in this study for strain accumulation and recrystallization (RXN) are labelled.
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Figure 2. Equilibrium phases as a function of temperature for the 8 wt% Ni steel.
Figure 2. Equilibrium phases as a function of temperature for the 8 wt% Ni steel.
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Figure 3. Optical microstructure of the as-cast Fe-21Mn-9Al-1C-8Ni after solution treatment at 1200 °C for 2 h and agitated water quenching. The matrix was primarily austenitic with interdendritic B2-NiAl.
Figure 3. Optical microstructure of the as-cast Fe-21Mn-9Al-1C-8Ni after solution treatment at 1200 °C for 2 h and agitated water quenching. The matrix was primarily austenitic with interdendritic B2-NiAl.
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Figure 4. (a) Optical micrographs and (b) secondary electron image of the Fe-21Mn-9Al-1C-8Ni steel hot rolled at 1050 °C with a finish rolling temperature of 900 °C followed by a 3 min precipitation harden at 1050 °C and water quenching (P1). Several non-recrystallized austenite grain boundaries contain coarsened B2-NiAl precipitation are delineated by dotted lines.
Figure 4. (a) Optical micrographs and (b) secondary electron image of the Fe-21Mn-9Al-1C-8Ni steel hot rolled at 1050 °C with a finish rolling temperature of 900 °C followed by a 3 min precipitation harden at 1050 °C and water quenching (P1). Several non-recrystallized austenite grain boundaries contain coarsened B2-NiAl precipitation are delineated by dotted lines.
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Figure 5. Optical micrographs of the Fe-21Mn-9Al-1C-8Ni alloy hot rolled at 1200 °C with a finish rolling temperature of 1050 °C followed by one cycle of 1050 °C hot rolling with a finish rolling temperature of 900 °C and water quench in the (a) as-rolled condition and (b) as-rolled and precipitation-hardened at 950 °C for 10 min (P2). (c) Secondary electron images of the P2 processed steel precipitation. Plate-type B2-NiAl were precipitated uniformly throughout the austenite matrix that nucleate and grow in two orientations within a grain.
Figure 5. Optical micrographs of the Fe-21Mn-9Al-1C-8Ni alloy hot rolled at 1200 °C with a finish rolling temperature of 1050 °C followed by one cycle of 1050 °C hot rolling with a finish rolling temperature of 900 °C and water quench in the (a) as-rolled condition and (b) as-rolled and precipitation-hardened at 950 °C for 10 min (P2). (c) Secondary electron images of the P2 processed steel precipitation. Plate-type B2-NiAl were precipitated uniformly throughout the austenite matrix that nucleate and grow in two orientations within a grain.
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Figure 6. Bright field STEM images of B2-NiAl precipitation within the austenite matrix in the P2 specimen with the corresponding SADP indicated by arrows for (a) the [013] zone axis of austenite and (b) the [011] zone axis for B2-NiAl.
Figure 6. Bright field STEM images of B2-NiAl precipitation within the austenite matrix in the P2 specimen with the corresponding SADP indicated by arrows for (a) the [013] zone axis of austenite and (b) the [011] zone axis for B2-NiAl.
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Figure 7. Tensile properties of the current 8% Ni steel in the P1 and P2 conditions, in comparison to a previously reported Fe-18.3Mn-10.1Al-0.9C-4.5Ni steel [42]. (a) The engineering and (b) true stress–strain curves, and (c) the work hardening rate as a function of true strain. (d) The log–log plots of true stress as a function of true plastic strain showing a non-linear behavior.
Figure 7. Tensile properties of the current 8% Ni steel in the P1 and P2 conditions, in comparison to a previously reported Fe-18.3Mn-10.1Al-0.9C-4.5Ni steel [42]. (a) The engineering and (b) true stress–strain curves, and (c) the work hardening rate as a function of true strain. (d) The log–log plots of true stress as a function of true plastic strain showing a non-linear behavior.
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Figure 8. Secondary electron image of the P2 condition in (a) the CVN fracture surface broken at −40 °C in the L-T orientation and (b) the tensile fracture surface. White arrows indicate regions of brittle fracture of austenite grains.
Figure 8. Secondary electron image of the P2 condition in (a) the CVN fracture surface broken at −40 °C in the L-T orientation and (b) the tensile fracture surface. White arrows indicate regions of brittle fracture of austenite grains.
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Figure 9. Equilibrium phases as a function of temperature for the 5 wt% Ni (dashed lines) and 8 wt% Ni (solid lines) steels. Arrows indicate the change in equilibrium phases from the 5 wt% Ni to 8 wt% Ni compositions.
Figure 9. Equilibrium phases as a function of temperature for the 5 wt% Ni (dashed lines) and 8 wt% Ni (solid lines) steels. Arrows indicate the change in equilibrium phases from the 5 wt% Ni to 8 wt% Ni compositions.
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Figure 10. Specific yield strength (SYS) as a function of ductility for various low-density Fe-Mn-Al-C steels in comparison to the currently studied P2 specimen. Shaded dots within circled regions represent previously reported Ni-containing Fe-Mn-Al-C steels [40,42,51,52,53,54,55,56,57,58].
Figure 10. Specific yield strength (SYS) as a function of ductility for various low-density Fe-Mn-Al-C steels in comparison to the currently studied P2 specimen. Shaded dots within circled regions represent previously reported Ni-containing Fe-Mn-Al-C steels [40,42,51,52,53,54,55,56,57,58].
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Table 1. Chemical analysis by OES and gas combustion analysis * in weight percent.
Table 1. Chemical analysis by OES and gas combustion analysis * in weight percent.
Chemistry (wt%)FeMnAlC *NiSiS *
Bal.21.48.91.037.70.050.0038
Table 2. Intragranular B2-NiAl dimensions and precipitate free zone (PFZ) width as a function of hot roll procedure.
Table 2. Intragranular B2-NiAl dimensions and precipitate free zone (PFZ) width as a function of hot roll procedure.
Specimen IDLength (nm)Width (nm)Aspect RatioPFZ Width (µm)
P1559 ± 240252 ± 1002.2N/A
P2124 ± 56* 7.8 ± 315.90.69 ± 0.13
* B2-NiAl thickness of the plate.
Table 3. Hardness, B2-NiAl morphology, and mechanical properties of the austenitic 8% Ni steel compared to the duplex 5% Ni steel as a function of alloying, thermomechanical processing, and heat treatment.
Table 3. Hardness, B2-NiAl morphology, and mechanical properties of the austenitic 8% Ni steel compared to the duplex 5% Ni steel as a function of alloying, thermomechanical processing, and heat treatment.
Steel & ProcessingHardness [HRC]B2-NiAl a MorphologyYS
[MPa]		UTS
[MPa]		% Total Elong.−40 °C Impact Toughness
L-T [J]		−40 °C Impact Toughness
T-L [J]		T-L/L-T
Ratio
P132.6 ± 1Types 2 and 3780 ± 321100 ± 1522.1 ± 5.137.2 ± 2.323.7 ± 0.50.64
P240.5 ± 1Types 2 and 3940 ± 71320 ± 526.6 ± 3.811.8 ± 1.0N/A bN/A
5Ni [42]32.7 ± 1Types 1, 2,
and 3		810 ± 421120 ± 12826.4 ± 1528.5 ± 2.7 c
(38.0 ± 3.6)		7.3 ± 0.2 c
(9.7 ± 0.3)		0.26
a Types 1, 2, and 3 refer to the B2-NiAl stringer, intergranular B2-NiAl, and intragranular B2-NiAl, respectively. b Not determined because of limited material. c 3/4 subsize bars. Values in parentheses indicate the linearly upscaled toughness to full-size specimens for comparison.
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MDPI and ACS Style

Piston, M.; Bartlett, L.; Limmer, K.R.; Field, D.M.; Hornbuckle, B.C. Formation of Heterogeneous Nucleation of B2-NiAl in Hot Rolled Fe-Mn-Al-C Plate: A Novel Composition and Processing Route for Lightweight High Strength Steel Containing Nickel. Metals 2024, 14, 1342. https://doi.org/10.3390/met14121342

AMA Style

Piston M, Bartlett L, Limmer KR, Field DM, Hornbuckle BC. Formation of Heterogeneous Nucleation of B2-NiAl in Hot Rolled Fe-Mn-Al-C Plate: A Novel Composition and Processing Route for Lightweight High Strength Steel Containing Nickel. Metals. 2024; 14(12):1342. https://doi.org/10.3390/met14121342

Chicago/Turabian Style

Piston, Michael, Laura Bartlett, Krista R. Limmer, Daniel M. Field, and Billy C. Hornbuckle. 2024. "Formation of Heterogeneous Nucleation of B2-NiAl in Hot Rolled Fe-Mn-Al-C Plate: A Novel Composition and Processing Route for Lightweight High Strength Steel Containing Nickel" Metals 14, no. 12: 1342. https://doi.org/10.3390/met14121342

APA Style

Piston, M., Bartlett, L., Limmer, K. R., Field, D. M., & Hornbuckle, B. C. (2024). Formation of Heterogeneous Nucleation of B2-NiAl in Hot Rolled Fe-Mn-Al-C Plate: A Novel Composition and Processing Route for Lightweight High Strength Steel Containing Nickel. Metals, 14(12), 1342. https://doi.org/10.3390/met14121342

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