Patent Application
20210262074
The present disclosure is generally related to steel alloys.
In ferritic steels (including tempered martensitic alloys), Cu precipitates form with a metastable BCC structure within the BCC ferritic matrix and then later transform into incoherent FCC particles [7]. These Cu precipitates act as powerful strengtheners, producing yield strength increases of ˜200 MPa [8,9]. The exact mechanism is not well understood, but has been attributed to moving dislocations transforming the metastable BCC structure of the Cu precipitate into one resembling FCC [8], modulus misfit strengthening between BCC Fe and BCC Cu [7], or dislocation pileup at the incoherent interface of the FCC Cu precipitates [7,8]. Intermetallic β-NiAl (B2) has also been used to strengthen ferritic steels, forming ordered coherent nano-precipitates that provide yield strength increases on the order of 100-500 MPa [10,11]. Combined, both Cu and β-NiAl exhibit an intriguing synergy in ferrite, whereby the formation of one phase both promotes the formation and accelerates the precipitation hardening kinetics of the other [9,12,13]. Together, these two precipitate phases can increase the yield strength of the alloys by 400-500 MPa with minimal loss in ductility.
The precipitation behavior of Cu and β-NiAl within Austenitic alloys is fundamentally different from that in ferritic systems, and is also not as well understood. FCC copper precipitates form coherently within the FCC Austenite matrix and are much less effective strengtheners than in ferrite [14,15]. The ordered BCC β-NiAl phase forms intragranularly as nano-scale platelets, on the order of 100-200 nm diameter and ˜20 nm thickness, with a Kurjumov-Sachs (K-S) orientation relationship with the FCC matrix [16]. These precipitates serve as potent strengtheners that can generate yield strengths in the range 1000-1300 MPa [16-18].
Ni and Mn are the primary substitutional elements for stabilizing the Austenite across a wide range of temperatures and strains [19]. Ni is more effective, but is roughly an order of magnitude more costly than Mn [20]. There are also different precipitation behaviors of β-NiAl in Ni-stabilized [16,21] and Mn-stabilized [18,22-24] Austenitic steels. The Mn-containing alloys appear to require rolling steps prior to ageing to precipitate the fine intragranular β-NiAl platelets along dislocations, whereas the Ni-stabilized alloys do not [16,22,25], suggesting a higher barrier to nucleation of the β-NiAl precipitates in the Mn-stabilized steels. While there is comparatively little published literature on combining Cu and β-NiAl precipitates to strengthen Austenitic steels, it has been reported that Cu additions refine the distribution of β-NiAl and thus increase the alloy strength [26]. This indicates a possible synergy between these two precipitate types that reduces the β-NiAl nucleation barrier, similar to what was reported in ferritic alloys.
Disclosed herein is an alloy comprising 7-30 wt. % manganese, 1-15 wt. % nickel, 1-10 wt. % aluminum, 1-8 wt. % copper, 0-15 wt. % chromium, 0-5 wt. % molybdenum, 0-3 wt. % vanadium, 0-3 wt. % titanium, 0-3 wt. % niobium, 0-2 wt. % silicon, 0-1 wt. % carbon, and balance of iron. A majority of the iron is γ-Fe. The alloy comprises β-NiAl precipitates, wherein at least 95 vol. % of the β-NiAl precipitates have a maximum dimension of 500 nm or less. The alloy comprises Cu-rich precipitates comprising at least 40 at. % copper.
Also disclosed herein is method comprising: providing a mixture of elements comprising 7-30 wt. % manganese, 1-15 wt. % nickel, 1-10 wt. % aluminum, 1-8 wt. % copper, 0-15 wt. % chromium, 0-5 wt. % molybdenum, 0-3 wt. % vanadium, 0-3 wt. % titanium, 0-3 wt. % niobium, 0-2 wt. % silicon, 0-1 wt. % carbon, and balance of iron; forming an alloy from the mixture; heating the alloy to a temperature that causes formation of γ-Fe; cooling or quenching the alloy to retain the γ-Fe at room temperature; and heat treating the alloy through one or more ageing steps to form β-NiAl precipitates and Cu-rich precipitates. At least 95 vol. % of the β-NiAl precipitates have a maximum dimension of 500 nm or less. The Cu-rich precipitates comprise at least 40 at. % copper.
A more complete appreciation will be readily obtained by reference to the following Description of the Example Embodiments and the accompanying drawings.
In the following description, for purposes of explanation and not limitation, specific details are set forth in order to provide a thorough understanding of the present disclosure. However, it will be apparent to one skilled in the art that the present subject matter may be practiced in other embodiments that depart from these specific details. In other instances, detailed descriptions of well-known methods and devices are omitted so as to not obscure the present disclosure with unnecessary detail.
To achieve an industry-scalable modern steel, the use of integrated computational materials engineering (ICME) [1] is essential for an agile alloy design process. The present study focuses on the application of ICME tools to design and develop a new family of fully Austenitic (γ-Fe FCC matrix), high-strength steels that exhibit the excellent ductility typically associated with Austenitic alloys. Commercially available Austenitic steels can only achieve yield strengths in the general range of σy=170-380 MPa [2-5]. Since solid-solution strengthening is much less effective in Austenitic steels than in ferritic steels (which have an a-Fe BCC matrix) [6,7], precipitation hardening was pursued as the primary strengthening mechanism.
Disclosed herein is an Austenitic steel (face centered cubic crystal structure) which is stabilized using a combination of C, Mn, Cu, and Ni and which is strengthened upon age heat treatment by the formation of a hierarchical structure of multiple precipitate phases, including intermetallic β-NiAl (B2 ordered body centered cubic crystal structure) and Cu precipitates, and optionally a carbide phase (for example the M23C6 (M=Mn, Cr, Mo) or the MC (M=V, Ti, Nb, Mo) carbides). These precipitates are achievable on the nano-scale for enhanced hardening, and can be tailored to obtain the desired mechanical properties and cost. The result is a high-strength Austenitic steel of lower-cost than Ni-based alternatives.
The steel (Fe-base alloy) is comprised of an Austenitic-structure matrix (face centered cubic crystal structure) stabilized by C, Ni, Cu, and Mn additions, with greater Mn content than Ni by weight for cost-effectiveness. Upon aging heat treatment (for example, times in the range of 1-24 hours at temperatures in the range 400-750° C.), multiple additional material phases precipitate, including a Cu-rich FCC phase, ordered intermetallic β-NiAl (B2 structure), and optionally carbide phase(s) (M23C6, MC). These precipitates are achievable on the nano-scale (on the order of 1-100 nm length) for increased material hardness. Because this nano-scale precipitation can be achieved without the use of mechanical processing steps prior to age heat treatment (e.g. cold rolling), further cost savings may be achieved. The discovery was made in a steel of approximate composition (percent by weight): Fe-17.7Mn-4.7Cr-0.48C-10Ni-5Al-4Cu. The alloy may contain the weight percentages of the elements in the Table 1, including any value within these ranges. The alloy may contain additional elements not listed in Table 1, with the balance of iron reduced accordingly.
The alloy may be made solely by thermal processing steps. Mechanical and thermomechanical processing, such as rolling, drawing, and peening, is optional. The elements are alloyed by any method for alloying metals, such as casting or arc-melting. The alloy is then heated to a temperature that converts the iron to γ-Fe. For example, the heating may be to 1000° C. under argon.
The sizes of the β-NiAl precipitates are such that at least 95 or 98 vol. % of the β-NiAl precipitates have a maximum dimension of 500 nm or less or 100 nm or less. As used herein “maximum dimension” is equivalent to a maximum caliper dimension and means the greatest length that can be measured for the precipitate between any two points on the precipitate surface. No length measurement of the precipitate is larger than this size no matter what direction that length is measured in. The vol. % of the precipitates refers to the vol. % of all of the β-NiAl precipitates in the alloy. None are excluded to make this calculation. Thus, very little, if any, of the β-NiAl precipitates are in the form of stringers, which can be detrimental to the mechanical properties of the steel [24]. The Cu-rich precipitates are at least 40 at. % copper.
The alloy may have a microhardness of at least 300 HV or 490 HV. The yield strength, which may be measured or estimated from the microhardness, may be at least 550 MPa, 689 MPa, or 1200 MPa.
The following examples are given to illustrate specific applications. These specific examples are not intended to limit the scope of the disclosure in this application.
Thermodynamic calculations of phase equilibria in potential alloys were conducted using the 2020a release of the Thermo-Calc software (Thermo-Calc, Stockholm, Sweden) with the TCFE9 thermodynamic database. To model the β-NiAl phase, the “B2_BCC” phase description included in TCFE9 was enabled, and a new composition set was defined with major sublattice constituents (Ni,Fe)(Al,Mn)(Va) due to high solubility for Fe and Mn in β-NiAl [27-29]. A baseline Austenitic composition of Fe-17.7Mn-4.7Cr-0.48C wt. % was selected from prior experimental studies [30,31]. Promising compositional modifications (Table 2) containing Ni+Al (Base-10-5-0), Cu (Base-0-0-4), and Ni+Al+Cu (Base-10-5-4) were identified with the guidance of Thermo-Calc modeling. These compositions were prepared by arc-melting high-purity (>99.9%) elemental constituents to produce ˜40 g button ingots in an ultra-high purity argon static atmosphere. The ingots were melted and inverted six times to ensure alloy homogeneity.
After melting, the ingots were sectioned, wrapped in Ta foil as an oxygen getter, and encapsulated in evacuated quartz ampoules backfilled with ˜34 kPa of UHP Ar. Samples were solutionized at 1000° C. for 12 h followed by 1100° C. for 96 h, which was calculated to be in the fully Austenitic region of the alloys, above the solvus of any carbides, Cu, or β-NiAl particles. The samples were water-quenched after the 1100° C. solutionizing treatment then re-encapsulated and aged at 580° C. for a variety of times, followed by water quenching.
After heat-treatment, the samples were sectioned, mounted in epoxy, and polished. Microhardness was carried out in a LECO AMH55 automatic hardness tester (LECO Corp., St. Joseph Mich.) with a Vickers-type indenter at 500 gf load and 15 second dwell time. Reported microhardnesses are the average of 20 measurements from each sample. Estimated values for yield strength (σy) were calculated from these experimental Vickers hardness values (HV) based on the empirical relation observed between these parameters in Austenitic alloys by previous authors [32].
The underlying nano-scale microstructures were analyzed by atom-probe tomography (APT) using a Cameca 4000× Si™ local electrode atom probe (LEAP) (Cameca, Gennevilliers Cedex, France) with a 355 nm ultraviolet pulsed laser, 40 K specimen base temperature, 30 pJ laser pulse energy, 500 kHz pulse repetition rate, and a detection rate of 0.05 ions per pulse (0.5%). Specimens for APT were prepared using standard lift-out and milling procedures [33-35] in a ThermoFisher Nova 600 NanoLab DualBeam™ focused ion beam/scanning electron microscope (FIB/SEM). Data reconstruction and analysis was done with Cameca Integrated Visualization and Analysis Software (IVAS) version 3.8.2. The β-NiAl, FCC—Cu, and M23C6 precipitates were delimited by isoconcentration surfaces of 20 at. % (Ni+Al), 5 at. % Cu, and 12 at. % Cr respectively, employing a voxel size of 1.0 nm, a delocalization distance of 3.0 nm laterally and 1.5 nm along the analysis direction, and a confidence sigma parameter of 1.0 for effective noise suppression [36]. These isoconcentration surface values were chosen to faithfully represent the extent of the precipitates and are consistent with those used in prior APT studies on similar alloys [9,15,37-43].
Results of the initial Thermo-Calc modeling are shown in
The ITIC phase diagram in
The Austenitic nature of the alloys was confirmed by X-ray diffraction (results not shown). Vickers microindentation was used as the primary indicator for the presence of strengthening precipitates due to their nano-scale.
The near-peak aged condition of the Base-10-5-4 alloy, 10 h at 580° C., was studied by APT and is displayed in
While Thermo-Calc predicts that ageing the Base-10-5-0 alloy will produce appreciable β-NiAl precipitation (
The barrier to β-NiAl precipitation can be overcome with the addition of Cu. A pronounced synergy was observed in the Base-10-5-4 alloy, where precipitates of β-NiAl, Cu, and M23C6 combine to increase the microhardness by 330 HV after ageing for 10 h at 580° C. This increase matches the observed improvement from β-NiAl precipitation in Ni-stabilized steels by other authors [16]. Because similar hardening was not observed in the Base-10-5-0 or Base-0-0-4 alloys, this demonstrates that Cu facilitates the nucleation of the β-NiAl precipitates to give rise to that hardening. This synergy is supported by APT analyses, which appear to show co-localization of the Cu and β-NiAl phases (
The Thermo-Calc predictions provide useful guidance for this ICME alloy design approach, but it is important to note the differences between those predictions and the experimental results to support improvements in the databases. One of the more important differences is in the stability of Austenite present at low temperatures. While
The apparent composition of the FCC—Cu precipitates, as measured by APT, is only ˜60 at. % Cu, much lower than that reported in other steel systems or suggested by the Fe—Cu binary phase diagram [15,53,54]. This is likely due to the well-documented local magnification effect in APT for Fe—Cu systems [55-57]. Thermo-Calc predicts this phase to dissolve at ˜450° C. in the Base-10-5-4 alloy, but Cu also incorporates into β-NiAl. The incorporation of Cu in β-NiAl is experimentally measured by APT to be ˜4 at. %, while Thermo-Calc predicts a much higher 9 at. % concentration (Table 3d). This discrepancy is consistent with the predicted lack of FCC—Cu precipitates in Base-10-5-4 at 580° C. (
Overall, this study has demonstrated the promise of combined Cu+β-NiAl nano-precipitates to strengthen Austenitic steels, for which a paucity of data presently exists [58]. The APT data of the near-peak aged condition presented here directly demonstrate the formation of a complex nano-structure that provides significant hardening to the Austenitic matrix upon ageing.
Obviously, many modifications and variations are possible in light of the above teachings. It is therefore to be understood that the claimed subject matter may be practiced otherwise than as specifically described. Any reference to claim elements in the singular, e.g., using the articles “a”, “an”, “the”, or “said” is not construed as limiting the element to the singular.
This application claims the benefit of U.S. Provisional Application No. 62/980,766, filed on Feb. 24, 2020. The provisional application and all other publications and patent documents referred to throughout this nonprovisional application are incorporated herein by reference.
| Number | Date | Country | |
|---|---|---|---|
| 62980766 | Feb 2020 | US |
