HIGH NITROGEN STEEL POWDER AND METHODS OF MAKING THE SAME

TECHNICAL FIELD

This disclosure relates to steel powders having dissolved nitrogen in excess of the solubility limit at solidification temperature and methods for making such powders with tailored phase constituents.

BACKGROUND

Nitrogen (N) is effective in improving the mechanical, wear and corrosion properties of steels if it remains in solid solution, specifically in the form of coherent Cr—N short range order (SRO). These steels are known as high nitrogen steels (HNS), comprising of significant amount of dissolved N (˜0.4-1.0 weight percent (wt %)), along with other alloying elements such as Cr and Mn and typically containing low nickel or no nickel. Austenitic stainless steels with dissolved nitrogen contents up to 0.60 wt % have successfully been utilized in applications involving pitting corrosion, crevice corrosion and stress corrosion cracking in hot chloride solutions, such as NaCl and MgCl₂. However, precipitation of nitride particles (e.g., Cr₂N, TiN, VN) leads to Cr and N depletion from the matrix, impairing the corrosion and wear resistance of the steel and should be avoided during processing.

The solubility of nitrogen in molten steels at atmospheric pressure is very low (0.045 wt % at 1600 degrees Celsius (° C.)). Normal steel making practice at atmospheric pressure does not permit dissolving high amounts of nitrogen into the melt and much of that dissolved nitrogen in the liquid is lost during liquid to δ ferrite solidification due to even lower N solubility in δ ferrite. One way to retain the dissolved nitrogen during liquid to solid transformation (or solidification) is by suppressing a liquid (L)→δ reaction and promoting a L→γ reaction due to high solubility of N in γ phase by alloying the steel with γ stabilizers such as Ni, Mn and N itself. The development of pressurized metallurgy, namely melting and solidifying under high N partial pressure, makes it possible to effectively enhance the dissolved N content without inducing nitrogen pores and utilize the beneficial role of dissolved nitrogen. But in general these procedures are very expensive and require sophisticated equipment.

Besides fabricating bulk high nitrogen steel (HNS) by pressurized metallurgy described above, powders of HNS can find many beneficial applications such as coatings and sintered metal products due to their excellent mechanical and corrosion properties. Steel powders are commonly fabricated by atomization of liquid steel. To produce HNS powder by an atomization technique, the melting procedure needs to use a high nitrogen partial pressure environment along with appropriate alloy composition to avoid an L→δ reaction and then atomizing the liquid with high pressure nitrogen gas jets. While this procedure is feasible, it is complicated and expensive. Alternatively, mechanical alloying by attrition milling can introduce high levels of nitrogen into powders; however, typical time requirements are in excess of 100 hours and only a limited amount of material can be processed at one time. Further, the process introduces undesirable impurities and hard powders. This is especially true for fabricating coatings using HNS powder, where there are not many good fabrication options available. These powders cannot be remelted using high temperature coating processes such as plasma spray as the dissolved nitrogen will be lost unless high nitrogen partial pressure is maintained during deposition process Cold spray processes that employ solid state fusion to fabricate coatings can be effective in this case, however, it requires that the powders possess sufficient deformability (plasticity) that necessitates well controlled nitrogen dissolution and phase constituents, preferably γ phase free of Cr₂N precipitates.

Providing a means to economically fabricate steel powders having high level of dissolved nitrogen without a nitride or oxide compound layer or damaging brittle nitride precipitates would benefit many industrial applications where a combination of high toughness, wear and corrosion properties are desirable.

SUMMARY

Provided are methods for the production of steel powders containing high dissolved nitrogen, in particular micron size un-sintered powders substantially free of brittle nitride compounds.

Accordingly, a method for producing high nitrogen steel (HNS) optionally in a powder form and from a powder precursor is provided, the precursor optionally comprising, ferrite (α) phase, or austenite (γ) phase or a mixture of α+γ phase; and the HNS powder comprising a mechanically tough alloy having dissolved nitrogen and optionally having a substantially homogeneous composition, in weight percent, of from 0.1 to 6.0 wt % nitrogen. Further, the HNS powder may optionally include a single phase nitrogen alloy, optionally a γ phase alloy. Thus, an HNS powder is provided comprising a dissolved nitrogen content significantly in excess of what could have been achieved through atomizing liquid steel at atmospheric pressures.

Further, an object of the present disclosure is to provide methodologies to remove preexisting oxides from the precursor steel powder to accelerate diffusion of nitrogen into the powder and prevent precipitation of incoherent nitride precipitates to promote plastic deformability. Methodologies as provided herein include exposing the precursor steel powder to a reducing gas environment at elevated temperatures and hydrogen from the reducing gas environment combines with the preexisting oxygen of the precursor, resulting in a volatile byproduct which is removed from the atmosphere, and further quenching the powders quickly to ambient temperatures after nitrogen dissolution to prevent precipitation of nitrides or formation of oxides. The oxygen removal methods optionally further include using a different gas composition from the ones used to introduce dissolved nitrogen or the same gas composition during the entire treatment cycle.

Also provided are methods for preventing sintering of powders during the process and promoting nitrogen uptake by the precursor powders. The methods as provided herein include continuously agitating the powder to prevent necking or joining between powders and thus maintaining supply route of nitrogen around the surface of the powder. The methods optionally include providing a rotary hot tube comprising baffles that prevent formation stratified layers and continuously breakdown any lumps formed. In other aspects, the methods include a fluidized bed reactor that uses the nitrogen containing gas and agitates, optionally continuously, the powder mass until the powder is quenched.

The above and other objects, features and advantages of the present disclosure will become more fully understood from the detailed description given herein below and the accompanying drawings which are given by way of illustration only, and thus are not to be considered as limiting.

BRIEF DESCRIPTION OF THE DRAWINGS

Exemplary aspects will become more fully understood from the detailed description and the accompanying drawings, wherein:

FIG. 1A is a schematic description of the solidification process of steel involving liquid to δ-ferrite transformation, followed by austenite and the associated rejection of nitrogen gas forming pores;

FIG. 1B is a schematic description of solidification process of steel involving liquid to austenite transformation and the associated retention of dissolved nitrogen gas in the solid precursor material according to the teachings of the current disclosure (exemplary aspect);

FIG. 2 is a schematic view of an exemplary microstructure of high nitrogen steel powder, having nitride precipitates;

FIG. 3 is an exemplary time-temperature cycle for fabricating HNS powder according to the some aspects of the current disclosure;

FIG. 4 is an exemplary outline of the inventive steps for fabricating HNS powder according to the teachings of the current disclosure;

FIG. 5 is a schematic composition map for adjusting the phase content in the HNS powder, according to the teachings of the current disclosure;

FIG. 6A is a schematic perspective view of an exemplary batch processing system for HNS powder wherein the processing chamber includes a rotary tube according to the teachings of the current disclosure;

FIG. 6B is a schematic cross sectional view of an exemplary batch processing system for HNS powder wherein the precursor powder is being loaded into the rotary tube by a screw feeder according to the teachings of the current disclosure;

FIG. 7A is a schematic cross sectional view of an exemplary batch processing system for HNS powder wherein the precursor powder is being processed within the rotary tube according to the teachings of the current disclosure;

FIG. 7B is a schematic cross sectional view of an exemplary batch processing system for HNS powder wherein the powder is being removed from the tube after nitrogen dissolution and quenched into a collection chamber according to the teachings of the current disclosure;

FIG. 8A is a schematic perspective view of an exemplary continuous processing system for HNS powder wherein the system includes a heated rotary tube and powder precursor is loaded at one end and the processed powder is collected at the opposite end according to the teachings of the current disclosure;

FIG. 8B is a schematic cross sectional view of an exemplary continuous processing system for HNS powder wherein the system includes a heated rotary tube and the powder precursor is loaded at one end and the processed powder is collected at the opposite end according to the teachings of the current disclosure;

FIG. 8C is a schematic cross sectional view of an exemplary continuous processing system for HNS powder wherein the system includes a heated rotary tube along with an auger and the powder precursor is loaded at one end and the processed powder is collected at the opposite end according to the teachings of the current disclosure;

FIG. 8D is a schematic cross sectional view of an exemplary continuous processing system for HNS powder wherein the system includes a heated rotary tube having a vibratory device and the powder precursor is loaded at one end and the processed powder is collected at the opposite end according to the teachings of the current disclosure;

FIG. 9 is a schematic perspective view of an exemplary fluidized bed powder processing system wherein the processing includes gas phase agitation according to some aspects of the teachings of the current disclosure;

FIG. 10 is a schematic cross sectional view of an exemplary fluidized bed powder processing system showing the precursor powder being loaded according to some aspects of the teachings of the current disclosure;

FIG. 11 is a schematic cross sectional view of an exemplary fluidized bed powder processing system showing the precursor powder being processed according to some aspects of the current disclosure;

FIG. 12 is a schematic cross sectional view of an exemplary fluidized bed powder processing system showing the processed powder being quenched according to the teachings of the current disclosure;

FIG. 13 is the schematic process map for powder processing showing different effects of temperature and holding time on the powder microstructure according to some aspects of the current disclosure:

FIG. 14 is the schematic cooling rates for powder quenching and their effects on the powder microstructure according to some aspects of the current disclosure;

FIG. 15 presents the Scanning Electron Microscope micrograph for powders processed under different process conditions according to some aspects of the current disclosure;

FIG. 16 presents the X-Ray diffraction patterns for as received solid precursor powder, and the powders processed under different conditions according to some aspects of the current disclosure.

DETAILED DESCRIPTION

Provided are processes forming high nitrogen steel and devices for performing the process. In particular, the processes as provided herein are useful for creating dissolved nitrogen in a solid steel material, optionally a powder material. The processes tailor heating and optionally holding periods at particular temperatures of the steel to form various phases, allow nitrogen to dissolve into the steel and prevent final formation of nitrides that hinder corrosion resistance and mechanical strength of the final high nitrogen steels.

The following terms or phrases used herein have the exemplary meanings listed below in connection with at least one embodiment:

“HNS” as used herein means steels having high nitrogen content specifically in dissolved solid solution form. The amount of nitrogen in the high nitrogen steel is optionally equal to or above the amount of nitrogen achievable in an equivalent steel alloy in a liquid state at atmospheric pressure of nitrogen.

“Precursor” as used herein means the starting steel powder used to make the HNS powder where the precursor powder has a lower nitrogen content that the resulting HNS powder.

“Compound” as used herein, means a material formed by reactions between elements having a stoichiometric ratio, illustratively, Cr₂N and Fe₂N, etc.

“Solid solution” as used herein, means an alloy formed by dissolving one or more alloying element(s) in a host solid without changing its phase. In specific aspects as provided herein, γ-Fe[N], wherein N is the alloying element dissolved in FCC-Fe, the austenite phase.

The addition of nitrogen improves the strength, ductility and impact toughness in austenitic steels, while the fracture strain and fracture toughness are not affected at elevated temperatures. The strength of nitrogen alloyed austenitic steels arises from three components: strength of the matrix, grain boundary hardening, and solid solution hardening. The matrix strength is not appreciably impacted by nitrogen, rather matrix strength correlates to the friction stress of the face centered cubic (FCC) lattice that is mainly controlled by the solid solution hardening of the substitutional elements like chromium and manganese. Grain boundary hardening, however, which occurs due to dislocation blocking at the grain boundaries, increases proportionally to the alloyed nitrogen content. The highest impact on the strength results from the interstitial solid solution of nitrogen. Nitrogen increases the concentration of free electrons promoting the covalent component of the interatomic bonding and the formation of Cr—N short range order (SRO). The occurrence of Cr—N SRO and the resultant interactions with dislocations and stacking faults are believed to play a major role in the deformation behavior of these alloys, and can be tailored to enhance the strength, ductility, and impact toughness.

The composition and temperature strongly influence the stacking fault energy (SFE) and in turn, the deformation mechanisms and strengthening behavior of austenitic steels. Increasing the SFE, causes the active deformation mechanisms to change and is generally favored to achieve pure dislocation glide and enhanced toughness. Specifically, the effect of N additions on the SFE in Cr and Mn alloyed steels is reported to be non-monotonic, exhibiting a minimum SFE at ˜0.4 wt % N. The decrease in SFE at low N content (e.g. less than 0.4 wt %) is believed due to the segregation of interstitial N atoms to stacking faults, however, at higher N contents (e.g. at or greater than 0.4 wt %) the SFE increases as the bulk effect of interstitial solid solution becomes more pronounced. However, the formation of nitrides such as Cr₂N, at elevated N content, affects the distribution of alloying elements within the lattice and in turn diminishes the bulk effect of interstitial solid solution and the SFE. The formation of nitrides occurs when the nitrogen content goes beyond certain threshold value (depends on the overall composition of the alloy) and should be discouraged to take advantage of the interstitial solid solution hardening phenomenon described above.

As such, the high nitrogen steels of the present disclosure are optionally free or substantially free of any nitrides. In steels containing alloying elements (e.g. Cr, Al, Mo, V, Ti, etc.) nitride formation occurs because these alloying elements are stronger nitride formers than iron. As such, nitrides of the type MxNy (where M is Cr, Al, Mo, V, Ti, etc., x any y are chosen to arrive at proper stoichiometry) develop with more propensity. The high nitrogen steels produced by the processes as provided herein are optionally absent, optionally substantially absent a nitride of Cr, Al, Mo, V, Ti, or others.

High nitrogen containing austenitic steels also exhibit excellent resistance to atmospheric corrosion. However, the corrosion resistance is also strongly influenced by the nitrogen content. At low N content, the formation of a phase (an intermetallic compound with Cr) at the grain boundaries as well as the formation of nitrides such as Cr₂N at high nitrogen content are detrimental to the corrosion resistance of these steels. Best corrosion resistance can be achieved if all nitrogen is in solid solution, i.e. no nitrides such as Cr₂N are precipitated. It can be summarized that an optimal combination of toughness and corrosion resistance can be achieved by limiting the nitrogen content within a range, wherein a substantially or completely precipitation free homogeneous microstructure with N in solid solution form can be obtained. This range of dissolved N depends on other alloying elements present in the alloy as well as the process thermal history as discussed herein.

One approach to obtain a homogeneous dissolved nitrogen content in a steel alloy, specifically in austenitic steel is to (i) dissolve the nitrogen into the alloy in liquid state and then (ii) solidify the alloy without losing the dissolved nitrogen during solidification. However, both the tasks have their own challenges. For example, the nitrogen solubility in liquid iron at atmospheric pressure is very low (0.045 wt % at 1600° C.). Nitrogen solubility in a liquid alloy increases by the square root of the partial pressure (Sievert's square root law). Hence, to introduce higher nitrogen into liquid iron/steel, melting should be done in a high pressure nitrogen environment. Nitrogen alloying in the molten state may be achieved by high pressure induction or electric arc furnaces, pressure electro slag remelting furnace (PERS), and plasma arc and high-pressure melting with hot isostatic processing (HIP), etc.

Further, it is also known that the addition of certain elements such as chromium, manganese vanadium, niobium, and titanium increases the nitrogen solubility, while addition of elements such as carbon, silicon, and nickel reduces the nitrogen solubility. Hence, in order to induce high nitrogen concentrations into the melt, chromium and manganese can be added and nickel should be avoided. Furthermore, in some aspects, elements such as vanadium, niobium, and titanium, are absent or present in insignificant amounts as they are powerful nitride formers.

While chromium addition significantly enhances nitrogen solubility in the melt, it is also a strong δ-ferrite stabilizer. As illustrated in FIG. 1A, δ-ferrite solidification in iron alloys is associated with a wide solubility gap and a sudden drop 12 of nitrogen solubility in the material. In other words, a melt containing dissolved nitrogen 13, will lose most of its nitrogen during δ-ferrite solidification even though the subsequent lower temperature austenite phase can dissolve a much higher amount of nitrogen, 11. It is important to note that in ferritic steels, enhancing the dissolved nitrogen content 14 in the liquid by alloying additions and performing the melting operation under high nitrogen pressure would not retain the dissolved nitrogen in the δ phase due to the associated loss during δ-ferrite solidification. This leads to the formation of interdendritic pores 18, which results in degraded material quality and the loss of nitrogen in the final material. Therefore, to retain the enhanced dissolved nitrogen achieved through high nitrogen pressure melting and alloying adjustment and transfer it to the solid austenitic material, the δ-ferrite solidification should be avoided. However, if the solidification operation is carried out under high nitrogen partial pressure, the pores can be suppressed increasing the N content to some extent 15, and importantly after the δ→γ transformation, substantial amount of nitrogen 16 can be dissolved in the γ phase; the extent of which depends on the holding temperature, pressure and time.

Now referring to FIG. 1B, in the absence of δ-ferrite solidification, wherein the liquid directly solidifies into austenitic material, much of the dissolved nitrogen 12′ in the liquid state will be retained in the mixture of austenite and the liquid 18′ and subsequently transfer into the solid austenite phase 13′. It is to be noted that the austenite phase can have a significant amount of dissolved nitrogen 11′ and in order to achieve the saturation level 11′ the liquid may contain higher dissolved nitrogen 14′ to start with, which can be achieved only by high pressure melting and alloying adjustment. Further, under high nitrogen partial pressure the austenite can pick up more nitrogen 16′ and depending upon the temperature and time of holding the nitrogen content can reach the solubility limit 11′. The elimination of δ-ferrite solidification step can be achieved by carefully adjusting the composition of the alloy. To this end, manganese addition plays an important role. While enhancing the nitrogen solubility in the melt, manganese also suppresses the formation of δ-ferrite during solidification. As discussed above, the significant enhancement of strength in nitrogen alloyed austenitic steel comes from the formation of Cr—N SRO. Additionally, Cr enhances the resistance against atmospheric corrosion and hence is an important alloying addition. Further, the effect of manganese on enhancing nitrogen solubility is known to be two times less than the effect of chromium. Hence, significantly higher amount of Mn compared to Cr may be present in order to provide equivalent nitrogen solubility, eliminate δ-ferrite formation as well as achieve enhanced toughness and corrosion resistance. Another way to promote austenitic solidification and avoid degassing of nitrogen is to add carbon, however, carbon contents >0.1 wt. % have negative influence on corrosion resistance and ductility of the material and hence should be avoided.

One main problem for the production of austenitic steels containing high manganese is the strong segregation behavior of manganese that leads to heterogenic microstructure; which is detrimental to the mechanical behavior as well as corrosion resistance. Further, as discussed above, precipitation of σ phase or nitrides such as Cr₂N 24 as shown in FIG. 2 should be avoided during processing to achieve high toughness and corrosion resistance. The segregation and precipitation issues can be suppressed or completely eliminated by rapidly cooling.

As a way of background, when austenitic steel is exposed to a nitrogen atmosphere at high temperature, nitrogen may be incorporated in steel through dissolution in the austenitic phase up to its solubility limit, according to equation (1):

½N₂(gas)⇔[N]γ (1)

and through precipitation of chromium nitrides, according to equations (2) and (3):

[Cr]γ+[N]γ⇔[CrN]γ (2)

2[Cr]γ+[N]γ⇔Cr₂N (3)

Nitrogen in Eqn (1) remains in solid solution depending on temperature of thermo-chemical treatment and nitrogen pressure. Nitrogen loss or nitrogen pickup may occur according to Sieverts' law at a given set of temperature and nitrogen partial pressure parameters. Also, note that CrN in Eqn. (2) is a coherent precipitate and is beneficial in enhancing mechanical properties of the steel, whereas Cr₂N is a precipitate that deteriorates the corrosion resistance. Further, due to slow diffusion rate of N in steel, nitrogen pick up is fast when the surface area is large which is achieved by exposing the powder to a nitrogen atmosphere. However, the presence of oxide layer on the steel surface inhibits the diffusion of nitrogen and should be removed for enhance the rate of diffusion. This can be achieved by treating the powder in a reducing gas atmosphere. Once the nitrogen is dissolved in the steel, it can be retained in the powder by quenching the powder to a temperature where the nitrogen diffusion is virtually absent.

Provided are methods for making nitrogen steel powder with a dissolved nitrogen content, the said dissolved nitrogen content optionally higher than the solubility limit of N in the alloy in its liquid state at atmospheric pressure and optionally the nitrogen alloy powder being devoid of a nitride compound precipitates or nitride compound layer. Referring to FIGS. 3 and 4, exemplary methods for the fabrication of the nitrogen alloy powder are provided. The temperature and time cycles 30 and method 40 may include one or more of the following steps; providing a solid precursor, optionally in powder form, surface form, coating form or other, the solid precursor steel with a substantially low dissolved nitrogen content in step 41 and the disposing the solid precursor powder into a nitrogen gas or mixture of a reducing gas and nitrogen gas environment in steps 42-44, where the said precursor powder can undergo an exemplary temperature-time cycle 32 or 33 to obtain the nitrogen steel powder with a dissolved nitrogen content in step 45. Further descriptions on steps 42-44 and temperature-time cycles 32 and 33 are provided below.

The reducing gas can optionally be a mixture of nitrogen and hydrogen, argon and hydrogen, or anhydrous ammonia. Under the reducing gas, the oxides layers will be removed and facilitate nitrogen introduction.

In some aspects, a solid precursor material is in the form of a coating on another substrate or other steel type. A coating optionally has a thickness, optionally from 10 nm to 100 micrometers.

In some aspects a precursor steel is in the form of a powder. The solid precursor powder material can optionally be obtained by atomizing a liquid steel alloy in atmospheric pressures.

Optionally, the powder is continuously agitated to provide contact with the gas as well as prevent sintering. Various methods for providing continuous agitation are described in this disclosure.

The precursor powder has a powder size. The precursor powder size is optionally between 5 and 250 micrometers (μm), is optionally between 5 μm and 150 μm, optionally between 10 μm and 75 μm. Powder size is defined as the size that is appropriately sieved through a desired sieve where powder below a certain size will pass through a first sieve and will have size that will be retained by a smaller second sieve. Choice of sieve size depends on the desired powder size.

The precursor steel is predominantly Fe (i.e. 50 wt % or greater Fe) and optionally includes one or more other elements that will promote FCC structure. For example, a precursor optionally includes Mn. Mn, when present, may be provided at a weight percent of 0 to 35. Optionally, the weight percent of Mn is less than 30. Optionally, the weight percent of Mn is 19-27. Optionally, the weight percent of Mn is 20-26. The presence of N in such alloys serves to promote and stabilize a desired FCC structure even when the amount of Mn or other FCC promoting metal is less than 20 weight percent. As such, the dissolved N and Mn optionally work in concert to promote austenitic structure to the protective layer metal alloy. Optionally, the precursor powder includes Ni, which also promotes austenitic structure. Ni, when present, may be provided at a weight percent of 0 to 20. Since Ni reduces the N solubility in the protective layer, the Ni is optionally between 0 wt % to 5 wt %. The precursor powder may optionally include C, that when present, may be provided at a weight percent of 0 to 0.2. While C improves N solubility, it also reduces the toughness of the resulting alloy. Optionally, the C is present in the precursor powder at 0 wt % to 0.1 wt %.

As mentioned earlier, the strengthening mechanism in nitrogen alloy steel emerges from the formation of Cr—N SRO and hence Cr is optionally included in the provided N alloy. However, Cr is a δ-ferrite promoter as well as ferrite stabilizer. In order to control the phase of the steel, the ferrite stabilizing effect of Cr may be countered by adjusting the amount of N and/or Mn, both of which serve as austenite stabilizers. The precursor may include one or more other metals. Optionally, a precursor may include molybdenum. Mo, when present, may be provided at a weight percent of 0 to 5. Optionally, a precursor may include aluminum. When present Al may be provided at 0.01 wt % to 10 wt %. Al is optionally present at or less than 10 wt %, optionally at or less than 8 wt %, optionally at or less than 6 wt %.

Now referring to FIG. 3, T_N34d represents a temperature where nitrogen uptake in steel occurs through nitride formation. The exact temperature and form of nitride depends on the steel composition. Tγ 34c represents the temperature at which the steel transforms into austenite of FCC form. Again, the temperature depends on the steel composition. T_γN34b represents the temperature at which all the nitride compounds dissolve and nitrogen in the steel exists in dissolved nitrogen form. Tm 34a represents the melting point of the steel. When the precursor in step 41 is subject to a temperature-time cycle 32 in steps 42-44, depending on the powder size and composition, first nitrogen uptake is expected to occur through nitride formation during the ramp-up phase 32′, which will dissolve during holding the powder above 34b. Above 34b, all the nitrogen uptake is expected to occur elemental N form. The extent of N uptake will depend on the ramp-up time plus the holding time above 34b. To retain the N in the steel in dissolved form and prevent reformation of nitrides, the steel needs to be quenched 32″ below 34d, quickly, which can be achieved by various methods. The quenching rate influences the final microstructure of the HNS. Under slow cooling the N may precipitate into carbide particles. The gas pressure, temperature and time are adjusted according to the desired dissolved nitrogen content in the final HNS of step 45 and the composition of the precursor.

Introducing dissolved N directly into a given steel powder without the formation of substantial nitrides following temperature-time cycle 32 (FIG. 3) may present some challenges. The desired holding temperature in this case is too close to the melting temperature, T_M34a of the alloy. When precursors are disposed to this temperature range, they tend to join together by a process known as sintering in the art. Although according to this disclosure the presence of sintering is not always undesirable, in many aspects sintering may hinder the use of any resulting powder in a coating application process, and hence undesirable in some applications. Further, sintering hinders compositional homogeneity of the powders, particularly in regards to N. Holding the temperature in this range for a short period without sintering, however, would limit the uptake of N. In other words, for a low target N uptake one can adopt temperature-time cycle 32 without any detrimental effects due to sintering. However, to avoid sintering but to achieve relatively increased nitrogen uptake, optionally, one can adopt alternative temperature-time process.

Referring again to FIG. 3, following temperature-time cycle 33, one can optionally hold the powder in the nitrogen environment just above the nitride formation temperature, T_Nfor a nitrogen uptake time to facilitate the uptake of significant amount of N. Optionally, a temperature for a nitrogen uptake time is at or above the T_Nand below the T_γ of the alloy. Further, holding the powder at this temperature prevents sintering especially in the presence of powder agitation is used during the nitrogen uptake time as will be illustrated below. However, holding the powder at a temperature between the Tγ and T_Nmay lead to the formation of undesirable nitrides. Accordingly, according to some aspects, the temperature-time cycle 33, comprises of a second heating step, where the temperature is raised above T_γNand held there for a brief nitride conversion time to decompose the nitrides, yielding the desired dissolved nitrogen content, prior to quenching the powder. The nitride conversion time above T_γNis substantially shorter compared to nitrogen uptake time and thus prevents sintering. It will be appreciated that T_N, Tγ, T_γNdepend on the steel composition, and further the holding times also depend on the powder composition and size. The temperature-time cycles 32 and 33 are exemplary illustrations and many variations can optionally be adopted to achieve sinter/nitride free powder with a desired dissolved N content.

In some aspects, a nitrogen uptake time is in excess of 1 second. Optionally, a nitrogen uptake time may be indefinite, but is more commonly 1 hour or less. For larger steel pieces or larger powder sizes the nitrogen uptake time may be adjusted upward. In particular aspects, a nitrogen uptake time is from 1 second to 15 minutes, optionally 1 second to 100 seconds, optionally 1 second to 60 seconds, optionally 10 seconds to 100 seconds, optionally 30 seconds to 70 seconds, optionally 50 seconds to 60 seconds. A hold time may be sufficient to fully heat the precursor to the desired temperature or may hold the precursor at that temperature for the nitrogen uptake time.

A nitrogen uptake time may be at a constant temperature or may be at a varying temperature. A varying temperature during a nitrogen uptake time may be at or between T_Nand T_γ. The temperature may fluctuate or remain substantially constant, optionally varying by 5° C. or less.

The precursor with the nitrogen uptaken into the material may then be subjected to a further heating step whereby the precursor is heated to a temperature near, but not at or above the Tm. Optionally, the second heating step heats the precursor to a second temperature that is above a T_γNof the precursor and below a melting temperature for the precursor powder. As illustrated above, above the T_γNof the precursor, any nitrides formed in the uptake step or otherwise present in the steel material are converted into dissolved nitrogen. This further increases the weight percent of dissolved nitrogen and prevents unwanted characteristics that occur due to the presence of nitrides in the final high nitrogen steel.

The increased temperature above the T_γNof the precursor is optionally held for a nitride conversion time. A nitride conversion time is optionally any time to allow all, substantially all or any desired amount of nitrides within the precursor to be converted to dissolved nitrogen. A nitride conversion time is optionally 1 hour or less. In some aspects, a nitrogen conversion time is a short as possible so as to both convert the nitrides to dissolved nitrogen but also to prevent sintering (in some aspects). As such, a nitride conversion time is optionally less than 1 hour, optionally less than 20 minutes, optionally, less than 10 minutes, optionally less than 5, 4, 3, 2, 1, min. It has been observed that some sintering may occur when using particular precursor steel in the form of a powder at 10 minutes. As such, when a powder precursor is used, the nitride conversion time is optionally 10 minutes or less, according to some non-limiting aspects.

Once the desired dissolved nitrogen content is achieved in step 43, the powder is quenched to a temperature where the diffusion is virtually absent in step 44 following an exemplary temperature-time cycle 33. The cooling rate is critical to avoid nitride formation or reformation. Optionally, the powder cooling rate is between 1° C./s and 100° C./s, optionally the powder cooling rate is between 5° C./s and 50° C./s, optionally the powder cooling rate is above 10° C./s.

The atmospheric pressures used in the processes optionally are not required to exceed 1 atm as, in many aspects, 1 atm is sufficient to radically increase the amount of dissolved nitrogen in the high nitrogen steel relative to prior processes. However, in other aspects the atmospheric pressure is optionally above 1 atm, optional 2 atm or greater, optionally 3 atm or greater, optionally 4 atm or greater. In some aspects, the atmospheric pressure is less than that typically required to dissolve nitrogen in liquid steel. As such, an atmospheric pressure is optionally less than 10 atm. optionally less than 9 atm, optionally less than 8 atm, optionally less than 7 atm, optionally less than 6 atm, optionally less than 5 atm.

The resulting high nitrogen steel is provided with an exceptionally low nitride content, optionally 0.01 wt % or lower. Optionally, the resulting HNS has a nitride content at or below 0.02 wt %, optionally 0.03 wt %, optionally 0.04 wt %, optionally 0.05 wt %, optionally 0.1 wt % or lower.

The resulting high nitrogen steel optionally has a dissolved nitrogen content of 0.05 wt % to 6.0 wt %, or higher or any value or range therebetween. Optionally the dissolved nitrogen content of the HNS is at ore greater than 0.1 wt %, optionally 0.5 wt %, optionally 1 wt %, optionally 2 wt %, optionally 3 wt %, optionally 4 wt %, optionally 5 wt %, optionally 6 wt %. In many aspects the amount of dissolved nitrogen exceeds the solubility limit of nitrogen in the alloy (alloy of otherwise identical composition) in a liquid state at atmospheric pressure.

In some aspects, the resulting high nitrogen steel includes a ferrite (α) phase, austenite (γ) phase, or a mixture of α+γ phase. In some aspects, the alloy is predominantly a single phase. A single phase may be a ferrite phase or a gamma phase. Optionally, an alloy is predominantly or entirely a single phase, optionally a γ phase. The HNS produced by the processes as provided herein is optionally predominantly FCC structure, optionally 90% or greater FCC structure. Optionally, the HNS produced by the processes as provided herein are absent BCC structure throughout the HNS.

As discussed above, some elements act as austenite stabilizers while others promote ferrite. Further, the extent of their influence also varies considerably. For example, N is almost 20 times more effective in stabilizing austenite compared to Mn. Similarly, Cr is almost two times more effective than Mo in stabilizing ferrite. Therefore, to predict the phases of the iron alloys of this disclosure, it is appropriate to use a nitrogen equivalent as a predictor of austenite/ferrite composition in a N alloyed protective layer as presented in this disclosure. For iron alloys primarily containing Mn. Cr, and N alloying elements, the N and Cr equivalents can be expressed as: N_eq=10 (wt. % N)+0.25 (wt. % Mn)−0.02 (wt. % Mn)²+0.00035 (wt. % Mn)³and Cr_eq=wt. % Cr, respectively. Note that should any other elements be present in appreciable amount, whether austenite stabilizer or ferrite stabilizer, N_eq and Cr_eq is modified appropriately. Further, there is a lot of controversy regarding weight factors for each element and often they are empirically determined from experiments. But, there is a general agreement that N and C are the two most impactful austenite stabilizers. Since addition of C beyond 0.1 wt % is detrimental to the toughness, primarily the influence of N and Mn is considered here for exemplary illustration of alloy compositions.

Accordingly, the target alloy composition impact on phase stability is illustrated in FIG. 5, wherein the phase boundary between 100% austenite and the mixture austenite+ferrite is separated by a line which can be expressed as N equivalent=A×Cr equivalent−B. Based on experimentations, A is ˜0.98 and B is ˜11.5 for Ni free Fe—Mn—Cr—N alloy. Accordingly, exemplary alloy compositions will lead to the following outcomes as presented in Table 1. The impact of Mn content in stabilizing the austenite decreases as the content increases. For example, keeping the nitrogen concentration at 0.5 wt %, an increment of Mn content from 15 wt % to 30 wt %, decreases the N-eq from 5.27 to 3.65. Further, N concentration is the most influential factor in stabilizing the austenite. For example, by changing the N concentration from 0.5 wt % in alloy #4 to 0.7 wt % in alloy #5, results in an austenitic alloy even though significant amount of Cr (20 wt %) is present in the alloy. However, care must be taken not to increase the N content significantly beyond the stability zone especially when high amount of Cr is present to prevent Cr₂N precipitation as illustrated in FIG. 2. Alternatively, Mn addition can counter the influence of Cr and contribute towards the stability of austenite. Optionally, the N kept between 0.4 wt. % and 0.9 wt. %, Mn is kept between 19-27 wt % and the Cr is kept between 10-18 wt. %, the rest being iron.

TABLE 1

Alloy
N
Mn
Cr
N_eq
Cr_eq

#
(wt %)
(wt %)
(wt %)
(wt %)
(wt %)
Phase

1
0.5
15
13
5.27
13
γ

2
0.5
20
13
4.6
13
γ

3
0.5
30
13
3.65
13
γ

4
0.5
20
20
4.6
20
γ + α

5
0.7
20
20
6.6
20
γ

An exemplary alloy containing 15 wt % Cr, 25 wt % Mn and 0.7 wt % N and the remainder Fe would form an austenite phase which is preferred in many applications. In some aspects, a N alloy is or includes 13-14 wt % Cr, 20-26 wt % Mn, and 0.4-0.6 wt % N with the remainder being Fe.

Accordingly, the exemplary method described in FIGS. 3 and 4, can be practiced in various embodiments as illustrated in FIGS. 6-12 as are described herein below.

According to the teaching of this disclosure, an exemplary embodiment 60 for batch processing of powder shown in FIGS. 6A and 6B, includes a ceramic processing tube 65 which is rotated using a motor-gear arrangement 67. One end of tube 65 is connected to a powder feeding apparatus, which comprises of motor 61, precursor inlet 62 to load powder and an auger 63. This feeding apparatus can be inserted or retracted using lead screw arrangement 69 to evenly distribute powder through the processing tube 65. The other end of the tube 65 is operably coupled to a cooled collection chamber 66, which can be cooled by chilled gas or other fluid. The powder treatment tube is heated to a desired temperature by a heater 64.

Further details of processing and removal of the precursor is shown in FIG. 7A and FIG. 7B respectively. To enhance processing of the precursor powder, fins 79 can be added to the processing tube 75. In order to prevent any oxidation of the powder, the internal environment of the processing tube is always kept under nitrogen atmosphere. The fins can be ceramic rods placed along wall of the processing tube 75 as shown in cross section 78. In order to remove the precursor, a tilt angle is applied to the rotating processing tube 75 through a jack arrangement 72 which in turn empties the treated powder into the collection chamber 76.

Another exemplary embodiment 80 for continuous processing of powder shown in FIGS. 8A and 8B, includes a ceramic processing tube 85 which is rotated using a motor-gear arrangement 87. One end of tube 85 is operably connected to a powder feeding apparatus 81, via precursor inlet 82 to continuously feed powder. The other end of the tube 85 is operably coupled to a cooled collection chamber 86, which can be cooled by chilled gas or other cooling fluids. The powder treatment tube is heated to a desired temperature by a heater 84. The tube assembly can be placed at different angles relative to gravity via an adjustment device 88. This angle as well as the rotation speed of the tube 85 determines the residence time of the precursor powder in the treatment tube. Accordingly, various temperature-time cycles presented in FIG. 3 can be adopted through this arrangement to continuously treat the powder to achieve a desired dissolved N content in the powder.

In another exemplary embodiment 80′, the auger 87′ extends from the powder reservoir 82′ till the delivery end of the processing tube 85′. Optionally, the processing tube remains stationary, while the auger continuously agitates the powder inside the processing tube. Optionally, the processing tube also rotates while treating the powder. The pitch and rotational speed of the auger 87′ controls the feed rate of the precursor as well as the dwell time in the processing tube 85′. The auger eventually pushes the precursor into the collection chamber 86′. In such a system, auger is made of high temperature compatible materials such as ceramic to be able to operate at high temperatures in the processing tube.

In yet another exemplary embodiment 80″, precursor feeding can be achieved using electromagnetic vibration 81″. The feed rate of the precursor is controlled by regulating the vibration frequency. The feed rate can be further increased by tilting the setup using a jack arrangement 83″. The precursor is introduced into the processing tube 85″ using commercially available powder feeder 82″ such as thermal spray powder feeder. Due to continuous vibration of the processing tube 85″, the precursor powders are agitated and moved through the processing tube into the collection chamber 86″.

Referring to FIG. 9, embodiment 90 provides a methodology to alloy the said precursor with N via a fluidized bed reactor. Embodiment 90 includes a heating element 92, a processing tube 93, a precursor intake 91, a vent for the exhaust gas 95, a downstream gate valve 98 and a cooled collection chamber 97, operably connected to the processing tube 93 via gate valve 98. As illustrated in the cross sectional view FIG. 10, a known quantity of precursor 99 is introduced into the processing tube 93 via the intake port 91, while keeping the gate valve 98 closed. Subsequently, preheated nitrogen gas is continuously injected through nozzle 94 which fluidizes the precursor as shown in FIG. 11. The exhaust gas leaves the processing chamber through the vent 95, which can be recycled back to the chamber nozzle 94. The combined heat input from the heating element 92 and the preheated fluidizing gas keeps the precursor powder temperature at a desired range. The temperature-time cycle to achieve a target dissolved nitrogen content in the steel powder is adopted according to the teachings of this disclosure as illustrated in embodiments 30 and 40. After the desired N concentration is achieved in the powder, as shown in FIG. 12, the downstream valve 98 is opened to release the treated powder 99 into the collection chamber 97 through the cooled channel 96. It is understood that the fluidized bed treatment method as described herein can be achieved by alternative embodiments following similar mechanical and physical principles.

Example

A precursor powder with composition of Fe, 12 wt % Cr and 20 wt % Mn was centrifugally atomized at Ervin Technologies, Tecumseh, Mich., USA and classified to yield a size distribution of 10 μm to 60 μm. The powder was then processed according to the teachings of the present disclosure by utilizing an embodiment illustrated FIG. 8C. Nitrogen uptake in the powder was determined using Leco TC436 combustion analyzer at NSL Analytical Services, Inc., Ohio, USA. The phase of the powder was determined using a Rigaku Miniflex X-ray Diffractometer (Cu Kα radiation λ˜1.5402 Å).

The precursor powder was subject to various temperature-time cycles under N2+5% H2 gas mixture environment. The observations are presented in FIGS. 13-16. As illustrated in FIG. 13, when the process temperature remained between T₁(<980° C.) and Tγ, for example 930° C. nitride precipitated were always present in the final powder as shown FIG. 15D. In the early stages of the treatment process, especially when the treatment temperature was closer to Tγ, both α and γ phases along with grain boundary precipitates were observed as shown in FIG. 15B. The corresponding x-ray diffraction pattern B shown in FIG. 16, indicates the presence of α and γ phases. It is to be noted that the incoming precursor powder had single phase (α or ferrite), as demonstrated by the diffraction pattern A in FIG. 16. With increasing treatment time, the overall nitride precipitate content increased and the matrix alloy transformation into γ phase. Longer holding time led to powder sintering. Processing temperatures, between Tγ and T_N, formed nitride precipitates in the α matrix. No changes were observed below the threshold temperature T_th. When the process temperature remained above T₂(>1080° C.), for example, 1130° C., although the N uptake was accelerated, the powders sintered quickly. The corresponding powder microstructure is shown in FIG. 15A. Temperature-time cycle within the shaded region as illustrated in FIG. 13, always resulted in nitride free single phase powder as shown in FIG. 15C and the corresponding X-ray diffraction pattern is shown FIG. 16A. It is to be understood that the boundaries of this operational regime is not limiting and is based on set of experiments, which is not completely exhaustive. It is to be noted that factors such as alloy composition, powder size and agitation during the treatment would affect the area of this zone. Powder cooling rate played a very important role as illustrated in FIG. 14. Cooling rate of 1° C./s always led to precipitate formation although the treated powder started with a single phase homogenous alloy. Cooling rate of 10° C./s ensured precipitate free single phase powder at room temperature, provided the treated powder started with a single phase homogenous alloy. Accordingly, a treatment zone as illustrated in FIG. 13, can be constructed to achieve single phase homogeneous dissolved nitrogen alloy according to the teachings of this disclosure. As shown in Table 2, processing at 980° C. and 1080° C. for a dwell time of 300 s, followed by quenching to room temperature at 10° C./s, yielded nitrogen content of 0.3 wt % and 0.48 wt % respectively, without any precipitates and uniform austenite phase. Further, sintering of powder was also absent in these experiments.

TABLE 2

Temperature
Processing
Nitrogen
Quench

° C.
time (s)
uptake % wt
rate ° C./s
Phase

980
300
0.3
10
γ

1080
300
0.48
10
γ

While particular embodiments have been illustrated and described herein, it should be understood that various other changes and modifications may be made without departing from the scope of the claimed subject matter. Moreover, although various aspects of the claimed subject matter have been described herein, such aspects need not be utilized in combination. It is therefore intended that the appended claims cover all such changes and modifications that are within the scope of the claimed subject matter.

Various modifications of the present invention, in addition to those shown and described herein, will be apparent to those skilled in the art of the above description. Such modifications are also intended to fall within the scope of the appended claims.

It is appreciated that all reagents are obtainable by sources known in the art unless otherwise specified.

The foregoing description is illustrative of particular embodiments of the invention, but is not meant to be a limitation upon the practice thereof.

Various modes for carrying out the present invention are disclosed herein; however, it is to be understood that the disclosed modes are merely exemplary of the invention that may be embodied in various and alternative forms. The figures are not necessarily to scale; some features may be exaggerated or minimized to show details of particular components. Therefore, specific structural and functional details disclosed herein are not to be interpreted as limiting, but merely as a representative basis for teaching one skilled in the art to variously employ the present invention.

Reference is made in detail to compositions, aspects and methods of the present disclosure. It is also to be understood that this disclosure is not limited to the specific aspects and methods described herein, as specific components and/or conditions may, of course, vary. Furthermore, the terminology used herein is used only for the purpose of describing particular aspects of the present disclosure and is not intended to be limiting in any way.

It must also be noted that, as used in the specification and the appended claims, the singular form “a,” “an,” and “the” comprise plural referents unless the context clearly indicates otherwise. For example, reference to a component in the singular is intended to comprise a plurality of components unless explicitly noted otherwise.

Throughout this description, where publications are referenced, the disclosures of these publications in their entireties are hereby incorporated by reference to more fully describe the state of the art to which this disclosure pertains.

HIGH NITROGEN STEEL POWDER AND METHODS OF MAKING THE SAME

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