Patent Application
20140190594
This application deals with new classes of advanced steel alloys which may be used for tubular product production. The new classes of advanced steel have unique chemistries and operable mechanisms leading to advanced mechanical properties.
Steels have been used by mankind for at least 3,000 years and are widely utilized in industry comprising over 80% by weight of all metallic alloys in industrial use. Existing steel technology is based on manipulating the eutectoid transformation. The first step is to heat up the alloy into the single phase region (austenite) and then cool or quench the steel at various cooling rates to form multiphase structures which are often combinations of ferrite, austenite, and cementite. Depending on the cooling rate of the steel at solidification or thermal treatment, a wide variety of characteristic microstructures (i.e. pearlite, bainite, and martensite) can be obtained with a wide range of properties. This manipulation of the eutectoid transformation has resulted in the wide variety of steels available nowadays.
Non-stainless steels may be understood herein to contain less than 10.5% of chromium and are typically represented by plain carbon steel which is by far the most widely used kind of steel. The properties of carbon steel depend primarily on the amount of carbon it contains. With very low carbon content (below 0.05% C), these steels are relatively ductile and have properties similar to pure iron. They cannot be modified by heat treatment. They are inexpensive, but engineering applications may be restricted to non-critical components and general paneling work.
Pearlite structure formation in most alloy steels requires less carbon than in ordinary carbon steels. The majority of these alloy steels is low carbon material and alloyed with a variety of elements in total amounts of between 1.0% and 50% by weight to improve its mechanical properties. Lowering the carbon content to the range of 0.10% to 0.30%, along with some reduction in alloying elements increases the weldability and formability of the steel while maintaining its strength. Such alloys are classed as a high-strength low-alloy steels (HSLA) exhibiting tensile strengths from 270 to 700 MPa.
Advanced High-Strength Steels (AHSS) steels may have tensile strengths greater than 700 MPa and include types such as martensitic steels (MS), dual phase (DP) steels, transformation induced plasticity (TRIP) steels, and complex phase (CP) steels. As the strength level increases, the ductility of the steel generally decreases. For example, low-strength steel (LSS), high-strength steel (HSS) and AHSS may indicate tensile elongations at levels of 25% to 55%, 10% to 45% and 4% to 30%, respectively.
Much higher strength (up to 2500 MPa) has been achieved in maraging steels which are carbon free iron-nickel alloys with additions of cobalt, molybdenum, titanium and aluminum. The term maraging is derived from the strengthening mechanism, which is transforming the alloy to martensite with subsequent age hardening. The common, non stainless grades of maraging steels contain 17% to 18% nickel, 8% to 12% cobalt, 3% to 5% molybdenum and 0.2% to 1.6% titanium. The relatively high price of maraging steels (they are several times more expensive than the high alloy tool steels produced by standard methods) significantly restricts their application in many areas (for example, automotive industry). They are highly sensitive to nonmetallic inclusions, which act as stress raisers and promote nucleation of voids and microcracks leading to a decrease in ductility and fracture toughness of the steel. To minimize the content of nonmetallic inclusions, the maraging steels are typically melted under vacuum resulting in high cost processing.
The present disclosure relates to a method for producing metallic alloys of selected elemental composition. The alloys may include: (1) Fe—Cr—Ni—B—Si alloys which optionally may include one or more of V, Zr, Mn, W, Ti, Mo, Nb, Al, Cu, V and C; (2) Fe—Ni—B—Si alloys which optionally include Cu or Mn; (3) Fe—Cr—B—Si alloys which optionally include Cu, C or Mn; and (4) Fe—B—Si—Mn alloys which optionally include Cu or C. The atomic ratios of the identified elements may add up to 100. Impurities may be present at levels up to 10 atomic percent.
The detailed description below may be better understood with reference to the accompanying FIGS. which are provided for illustrative purposes and are not to be considered as limiting any aspect of this invention.
a illustrates additional production methods of using powder metallurgy to produce seamless tubular products.
The alloys described herein may be employed for use in tubular products for different applications, such as pipes, tubes, casings, and rods with different profiles. They can be directly cast into tubular products by centrifugal casting or used in a powder form for further production steps including but not limited to cold/hot extrusion, pressing, forging towards final products. The alloys herein in powder form can be used as initial precursor or may be pre-compacted into billets (preforms).
Tubular products for drilling applications may include use as a drill collars (a component that provides weight on a bit for drilling), drill pipe (hollow wall pipe used on drilling rigs to facilitate drilling), tool joints (i.e. the threaded ends of drill pipe), protective well casings, and wellheads (i.e. the component of a surface or an oil or gas well that provides the structural and pressure-containing interface for drilling and production equipment) including but not limited to ultra-deep and ultra-deep water and extended reach (ERD) well exploration. The alloys herein may also be produced in a tubular form for the automotive industry, for example, as air bag tubes.
A number of production methods are envisioned for seamless tubular products. Centrifugal casting can be used to produce a pipe form directly from a liquid melt. The as-cast pipe can be post processed by various methods including hot extrusion, cold extrusion, hot pilgering, and/or cold pilgering in order to eliminate defects, change metallurgical structure, reduce surface or internal porosity, and increase pipe uniformity and quality. Another route is to utilize a powder metallurgy compact which involves atomization by several methods including gas, water, or centrifugal atomization etc. followed by compacting into a billet. Compacting can be done by several methods including powder extrusion, cold isostatic processing (CIP), hot isostatic pressing (HIP), etc. The compacts can then be processed through a number of different post processing strategies to yield a tubular form including piercing, pierce and roll, hot extrusion, cold extrusion, hot pilgering, cold pilgering, etc. The tubes produced are envisioned to be seamless but can be welded consistent with the good weld properties of the Class 1, Class 2, and Class 3 steel characteristics. However, the greatest usage of the new steel classes would be envisioned to be in the production of seamless tubes which have homogenous walls without any weld or joint along its length. The homogenous wall and smooth inner surface of the tube is not subjected to weaknesses caused by welding and the potential deleterious effects of the heat affected zone. The largest benefit to using seamless steel pipe is the increased pressure ratings that significantly broadening up the areas of application for ultra-deep and ultra-deep water and extended reach (ERD) well exploration. Seamless pipes are also in high demand for industrial boilers including power industry. Seamless tubes find application in the manufacture of bearings, automobile parts, drill pipes, hydraulic cylinders, gas cylinders, etc.
Centrifugal casting is a commercially available manufacturing process for producing a wide variety of cylindrically symmetrical parts from simple shapes such as pipes, tubes and tubular components to complex shapes such as valve balls or flanges. The process consists of pouring a molten metal liquid into a rapidly rotating cylindrical mold. The centripetal acceleration of the cylindrical mold causes the metal liquid to be pushed radially outward against the inner surface of the mold. The apparent centrifugal force acting on the metal liquid is proportional to the radius of rotation and to the square of the rotational speed along with the mass. Thus the pressure applied to the liquid can be significant because of the rapid rotation and larger diameter molds will have a greater pressure exerted on the liquid. The pressure can be engineered for the process in order for the produced part to be fully dense and thus defect free. This gives centrifugal casting an advantage over other casting techniques in which porosity can routinely occur and cannot be avoided. In addition any metal oxide material or slag entrained in the molten liquid will be separated from the metal liquid because of the oxide material or slag has lower density than the metal alloy and these impurities will “float” to the center of the rotation where there is minimal pressure. If the casting shape is a tube or pipe then any oxide material will segregate to the inner diameter surface of the tube or pipe where it can be easily removed by post casting boring or machining.
The mold is water quenched on the outside, which causes the molten metal to be rapidly solidified on the inner diameter of the mold where the pressure is the greatest because it is at the maximum radius and the solidification propagates radially towards the center axis of rotation. The result is that there is refined grain size microstructure on the outer surface, which will provide higher strength in the material. Also because the mold is spun rapidly until solidification is complete therefore the pressure is continuously applied and there is no possibility of any shrinkage during the solidification process, which occurs in other casting techniques.
The design of the part to be cast determines the amount of molten liquid metal to be poured into the mold. For pipes or tubes the molten metal volume is less than the total volume of the mold and is carefully regulated in the manufacturing process to ensure part reproducibility. There is therefore insufficient molten metal to fill up the mold completely. Consequently, a symmetrical bore hole forms in the center along the axis of symmetry. The amount of liquid metal controls the diameter of the bore hole. For low volumes the bore hole will have a large diameter meaning that the part has a thin wall. Thus the same mold can produce different thicknesses simply by changing the volume of metal liquid used to cast the pipe, which means that a range of pipe thicknesses can easily be produced with the same mold. Centrifugal casting production begins with weighing out commercial grade feedstock constituents that are then alloyed together in a furnace. The molten liquid metal is then transferred to a tundish that is moved to the casting location. The mold is rotated at a fixed speed and the tundish is poured into a spout that feeds the liquid metal into one end of the mold. The liquid metal flows along the mold because of the hydrostatic pressure of the fluid reservoir in the spout while simultaneously fans out on the inner diameter surface of the mold because of the rotation of the mold. Once the liquid metal has flowed to the desired length cooling water is applied to the outside of the mold causing the molten metal to solidify. Centrifugal casting is conducted with either vertical or horizontal oriented molds. The choice of orientation used in the manufacturing process is mainly dictated by the length of the mold along its axis of symmetry. For short mold lengths, the mold is vertical while for tubes or pipes for which the length of the axis of symmetry is significantly longer than the diameter then the mold is oriented horizontally. A major driving factor for mold orientation is the mechanical method used to rotate the mold with most casting done with horizontal molds. This allows for mechanical drive wheels to be contact on the outside of the mold in order to precisely control the rotational speed. A centrifugal casting machine is schematically presented in
Conventional processing of pipe would involve continuous casting into thick rods which will be cut as preforms or billets to be used as feedstock for the pipe making process. Due to the unique metallurgy including the chemistry, operable mechanisms and targeted structures, the Class 1, 2, and 3 steels described in this application cannot be made by continuous casting of thick blocks or rods. However, there are a number of methods envisioned to produce pipe using a powder metallurgy process and one of the methods is described in
Extrusion is a process used to create objects of a fixed, cross-sectional profile and is widely used for seamless pipe and tube production. The main goal in tube extrusion is to manufacture consistent products with minimal dimensional variation. A material is pushed or drawn through a die of the desired cross-section. Extrusion is in most cases a hot working operation but can also be carried out in cold mode. The working temperatures in hot extrusion are typically 0.7-0.9 Tm, where Tm is the melting temperature. In steel extrusion, glass lubrication is commonly used when a layer of glass between the billet and the container, between the billet and the mandrel, and between the billet and the die is applied. Each billet is heated to the extrusion temperature and then rolled in a powder of glass during transportation to the extrusion chamber. Glass powder is also applied inside the billet to assure good lubrication between billet and mandrel. Lubrication through the die is provided by a thick disc of compacted glass, the glass pad, which is placed between the billet and the die. During extrusion, the glass pad is pressed against the die by the hot metal. The glass pad will deform with the billet and melt progressively to surround the extrusion with a lubricant glass film. The billet heating is an important stage in the manufacturing of steel tubes and pipes. The aim is to heat the material to a specified temperature that is suitable for hot forming. It is often desired to have a uniform temperature distribution within the heated billet. Heating prior to extrusion is carried out in gas-fueled rotary hearth furnaces, in induction furnaces, or in a combination of both.
a describes additional alternate powder metallurgy approaches to produce seamless pipe. After billet preparation, the billet can be extruded, pilgered or processed through a pierce-and-roll process. The extrusion process was described previously and can be done hot or cold. In the pilgering process, the preform billet, which can be heated up to a target process temperature, is forced through a die which generally results in reduction of both the outside and inner diameters. In the pierce and roll process, seamless pipe is produced in a typical four step process. The heated ingot/billet is first pierced to create room for the insertion of the mandrel and pipe is produced through what has been termed the Mannesman effect. Next a mandrel is inserted which maintains the inner diameter of the tube while it goes through the mandrel mill, which forms the outer diameter through a series of perpendicular orientated pairs of rolling mills. Next the pipe is further reheated and then goes through a stretching/finishing mill where pairs of rollers offset by 120 degrees are utilized to achieve final finishing/tolerances. Afterward, the pipe is often heat treated by various processes to hit the targeted structures and final properties and then is straightened.
The alloys herein are such that they are capable of formation of what is described as Class 1, Class 2 Steel or Class 3 Steel which are preferably crystalline (non-glassy) with identifiable crystalline grain size morphology. The ability of the alloys to form Class 1, Class 2 or Class 3 Steels herein is described in detail. However, it is useful to first consider a description of the general features of Class 1, Class 2 and Class 3 Steels, which is initially provided below.
The formation of Class 1 Steel herein is illustrated in
The Modal Structure of Class 1 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 500 nm to 20,000 nm containing austenite and/or ferrite; (2) boride grain size of 25 nm to 500 nm (i.e. non-metallic grains such as M2B where M is the metal and is covalently bonded to B). The boride grains may also preferably be “pinning” type phases which is reference to the feature that the matrix grains will effectively be stabilized by the pinning phases which resist coarsening at elevated temperature. Note that the metal boride grains have been identified as exhibiting the M2B stoichiometry but other stoichiometries are possible and may provide pinning including M3B, MB (M1B1), M23B6, and M7B3.
The Modal Structure of Class 1 Steel may be subjected to thermomechanical deformation and/or heat treatment, resulting in some variation in properties, but the Modal Structure may be maintained.
When the Class 1 Steel noted above is exposed to a mechanical stress, the observed stress versus strain diagram is illustrated in
Reference to the hexagonal phases may be understood as a dihexagonal pyramidal class hexagonal phase with a P63mc space group (#186) and/or a ditrigonal dipyramidal class with a hexagonal P6bar2C space group (#190). In addition, the mechanical properties of such second type structure of the Class 1 Steel are such that the tensile strength is observed to fall in the range of 700 MPa to 1400 MPa, with an elongation of 10-50%. Furthermore, the second type structure of the Class 1 Steel is such that it exhibits a strain hardening coefficient from 0.1 to 0.4 that is nearly flat after undergoing the indicated yield. The strain hardening coefficient is reference to the n-value in the formula σ=Kεn, where σ represents the applied stress on the material, ε is the strain and K is the strength coefficient. The value of the strain hardening exponent n lies between 0 and 1. A value of 0 means that the alloy is a perfectly plastic solid (i.e. the material undergoes non-reversible changes to applied force), while a value of 1 represents a 100% elastic solid (i.e. the material undergoes reversible changes to an applied force).
Table 1A below provides a comparison and performance summary for Class 1 Steel herein.
The formation of Class 2 Steel herein is illustrated in
As shown therein, Structure #1 is initially formed in which Modal Structure is the result of starting with a liquid melt of the alloy and solidifying by cooling, which provides nucleation and growth of particular phases having particular grain sizes. Grain size herein may again be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Accordingly, Structure #1 of the Class 2 Steel may be preferably achieved by processing through either laboratory scale procedures and/or through industrial scale methods such as powder atomization or alloy casting.
The Modal Structure of Class 2 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 500 nm to 20,000 nm containing austenite and/or ferrite; (2) boride grain size of 25 nm to 500 nm (i.e. non-metallic grains such as M2B where M is the metal and is covalently bonded to B). The boride grains may also preferably be “pinning” type phases which are referenced to the feature that the matrix grains will effectively be stabilized by the pinning phases which resist coarsening at elevated temperature. Note that the metal boride grains have been identified as exhibiting the M2B stoichiometry but other stoichiometries are possible and may provide pinning including M3B, MB (M1B1), M23B6, and M7B3 and which are unaffected by Mechanisms #1 or #2 noted above). Reference to grain size is again to be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Furthermore, Structure #1 of Class 2 steel herein includes austenite and/or ferrite along with such boride phases.
In
Characteristic of the Static Nanophase Refinement Mechanism #1 in Class 2 steel, the micron scale austenite phase (gamma-Fe) which was noted as falling in the range of 500 nm to 20,000 nm is partially or completely transformed into new phases (e.g. ferrite or alpha-Fe). The volume fraction of ferrite (alpha-iron) initially present in the Modal Structure (Structure 1) of Class 2 steel is 0 to 45%. The volume fraction of ferrite (alpha-iron) in Structure #2 as a result of Static Nanophase Refinement Mechanism #2 is typically from 20 to 80%. The static transformation preferably occurs during elevated temperature heat treatment and thus involves a unique refinement mechanism since grain coarsening rather than grain refinement is the conventional material response at elevated temperature.
Accordingly, grain coarsening does not occur with the alloys of Class 2 Steel herein during the Static Nanophase Refinement mechanism. Structure #2 is uniquely able to transform to Structure #3 during Dynamic Nanophase Strengthening and as a result Structure #3 is formed and indicates tensile strength values in the range from 800 to 1800 MPa with 5 to 65% total elongation.
Depending on alloy chemistries, nano-scale precipitates can form during Static Nanophase Refinement and the subsequent thermal process in some of the non-stainless high-strength steels. The nano-precipitates are in the range of 1 nm to 200 nm, with the majority (>50%) of these phases 10˜20 nm in size, which are much smaller than the boride pinning phase formed in Structure #1 for retarding matrix grain coarsening. Also, during Static Nanophase Refinement, the boride grain sizes grow larger to a range from 200 to 2500 nm in size.
Expanding upon the above, in the case of the alloys herein that provide Class 2 Steel, when such alloys exceed their yield point, plastic deformation at constant stress occurs followed by a dynamic phase transformation leading toward the creation of Structure #3. More specifically, after enough strain is induced, an inflection point occurs where the slope of the stress versus strain curve changes and increases (
With further straining during Dynamic Nanophase Strengthening, the strength continues to increase but with a gradual decrease in strain hardening coefficient value up to nearly failure. Some strain softening occurs but only near the breaking point which may be due to reductions in localized cross sectional area at necking. Note that the strengthening transformation that occurs at the material straining under the stress generally defines Mechanism #2 as a dynamic process, leading to Structure #3. By dynamic, it is meant that the process may occur through the application of a stress which exceeds the yield point of the material. The tensile properties that can be achieved for alloys that achieve Structure 3 include tensile strength values in the range from 800 to 1800 MPa and 5 to 65% total elongation. The level of tensile properties achieved is also dependent on the amount of transformation occurring as the strain increases corresponding to the characteristic stress strain curve for a Class 2 steel.
Thus, depending on the level of transformation, tunable yield strength may also now be developed in Class 2 Steel herein depending on the level of deformation and in Structure #3 the yield strength can ultimately vary from 400 MPa to 1700 MPa. That is, conventional steels outside the scope of the alloys here exhibit only relatively low levels of strain hardening, thus their yield strengths can be varied only over small ranges (e.g., 100 to 200 MPa) depending on the prior deformation history. In Class 2 steels herein, the yield strength can be varied over a wide range (e.g. 400 to 1700 MPa) as applied to Structure #2 transformation into Structure #3, allowing tunable variations to enable both the designer and end users in a variety of applications, and utilize Structure #3 in various applications such as crash management in automobile body structures.
With regards to this dynamic mechanism shown in
Note that dynamic recrystallization is a known process but differs from Mechanism #2 (
Class 3 steel is associated with formation of a High Strength Lamellae Nanomodal Structure through a multi-step process as now described herein.
In order to achieve a tensile response involving high strength with adequate ductility in non-stainless carbon-free steel alloys, a preferred seven-step process is now disclosed and shown in
Structure #1 involving the formation of the Modal Structures (i.e. bi, tri, and higher order) may be achieved in the alloys with the referenced chemistries in this application by processing through the laboratory scale as shown and/or through industrial scale methods involving chill surface processing such as twin roll casting or thin slab casting. The Modal Structure of Class 3 Steel will therefore initially indicate, when cooled from the melt, the following grain sizes: (1) matrix grain size of 500 nm to 20,000 nm containing ferrite or alpha-Fe (required) and optionally austenite or gamma-Fe; and (2) boride grain size of 100 nm to 2500 nm (i.e. non-metallic grains such as M2B where M is the metal and is covalently bonded to B); (3) yield strengths of 350 to 1000 MPa; (4) tensile strengths of 200 to 1200 MPa; and total elongation of 0-3.0%. It will also indicate dendritic growth morphology of the matrix grains. The boride grains may also preferably be “pinning” type phases which is reference to the feature that the matrix grains will effectively be stabilized by the pinning phases which resist coarsening at elevated temperature. Note that the metal boride grains have been identified as exhibiting the M2B stoichiometry but other stoichiometries are possible and may provide pinning including M3B, MB (M1B1), M23B6, and M7B3 and which are unaffected by Mechanism #1, #2 or #3 noted above). Reference to grain size is again to be understood as the size of a single crystal of a specific particular phase preferably identifiable by methods such as scanning electron microscopy or transmission electron microscopy. Accordingly, Structure #1 of Class 3 steel herein includes ferrite along with such boride phases.
Structure #2 involves the formation of the Modal Lath Phase Structure with uniformly distributed precipitates from Modal Structure (Structure 1) with dendritic morphology though Mechanism #1. Lath phase structure may be generally understood as a structure composed from plate-shaped crystal grains. Reference to “dendritic morphology” may be understood as tree-like and reference to “plate shaped” may be understood as sheet like. Lath structure formation preferably occurs at elevated temperature (e.g. at temperatures of 700° C. to 1200° C.) through plate-like crystal grain formation with: (1) lath structural grain sizes typically from 100 to 10,000 nm; (2) boride grain size of 100 nm to 2,500 nm; (3) yield strengths of 350 MPa to 1400 MPa; (4) tensile strengths of 350 MPa to 1600 MPa; (5) elongation of 0-12%. Structure #2 also contains alpha-Fe and gamma-Fe remains optional.
A second phase of boride precipitates with a size typically from 100 to 1000 nm may be found distributed in the lath matrix as isolated particles. The second phase of boride precipitates may be understood as non-metallic grains of different stoichiometry (M2B, M3B, MB (M1B1), M23B6, and M7B3) where M is the metal and is covalently bonded to Boron. These boride precipitates are distinguished from the boride grains in Structure #1 with little or no change in size.
Structure #3 (Lamellae Nanomodal Structure) involves the formation of the lamellae morphology as a result of static transformation of ferrite into one or several phases through Mechanism #2 identified as Lamellae Nanophase Creation. Static transformation is a decomposition of the parent phase into new phase or several new phases due to alloying elements distribution by diffusion during elevated temperature heat treatment, which may preferably occur in the temperature range from 700° C. to 1200° C. Lamellae (or layered) structure is composed of alternating layers of two phases whereby individual lamellae exist within a colony connected in three dimensions. A schematic illustration of lamellae structure is shown in
Lamellae Nanomodal Structure (Structure #3) transforms into Structure #4 through Dynamic Nanophase Strengthening (Mechanism #3, exposure to mechanical stress) during plastic deformation (i.e. exceeding the yield stress for the material) displaying relatively high tensile strengths in the range of 1000 MPa to 2000 MPa. In
The strengthening during deformation is related to phase transformation that occurs as the material strains under stress and defines Mechanism #3 as a dynamic process. For the alloy to display high strength at the level described in this application, lamellae structure is preferably formed prior to deformation. Specific to this mechanism, the micron scale austenite phase is transformed into new phases with reductions in microstructural feature scales generally down to the nanoscale regime. Some fraction of austenite may initially form in some Class 3 alloys during casting and then may remain present in Structure #1 and Structure #2. During straining when stress is applied, new or additional phases are formed with nanograins typically in a range from 1 to 100 nm.
In the post-deformed Structure #4 (High Strength Lamellae Nanomodal Structure), the ferrite grains contain alternating layers with nanostructure composed from new phases formed during deformation. Depending on the specific chemistry and the stability of the austenite, some austenite may be additionally present. In contrast with layers in Structure #3 where each layer represents a single or just few grains, in Structure #4, a large number of nanograins of different phases are present as a result of Dynamic Nanophase Strengthening. Since nanoscale phase formation occurs during alloy deformation, it represents a stress induced transformation and defined as a dynamic process. Nanoscale phase precipitations during deformation are responsible for extensive strain hardening of the alloys. The dynamic transformation can occur partially or completely and results in the formation of a microstructure with novel nanoscale/near nanoscale phases specified as High Strength Lamellae Nanomodal Structure (Structure #4) that provides high strength in the material. Thus the Structure #4 can be formed with various levels of strengthening depending on specific chemistry and the amount of strengthening achieved by Mechanism #3. Table 1C below provides a comparison of the structure and performance features of Class 3 Steel herein.
The formation of Modal Structure (MS) in either Class 1, Class 2 or Class 3 Steel herein can be made to occur at various stages of the production process. Accordingly, the formation of MS may depend specifically on the solidification sequence and thermal cycles (i.e. temperatures and times) that the alloy is exposed to during the production process. The MS may be preferably formed by heating the alloys herein at temperatures in the range of above their melting point and in a range of 1100° C. to 2000° C. and cooling below the melting temperature of the alloy, which corresponds to preferably cooling in the range of 11×103 to 4×10−2 K/s.
Subsequent thermal cycles and/or deformation at production including but not limited to HIP consolidation of powder into billets, hot extrusion, hot pressing, forging, as well as, post-production heat treatment will lead to Modal Nanophase structure formation in Class 1 alloys, Nanomodal or High Strength Nanomodal Structure formation in Class 2 alloys, and Lamellae Nanomodal or High Strength Lamellae Nanomodal Structure in Class 3 alloys. Depending on the Class of the alloy, its chemical composition and type of the microstructure formed at the production cycle, tubular products with wide variety of properties can be produced.
The chemical composition of the alloys studied is shown in Table 2 which provides the preferred atomic ratios utilized. In Case Examples below, the processability of the alloys in Table 2 by different technological methods and advanced property combinations can be achieved in the tubular products.
From the above, it can be seen that the targeted composition ranges (atomic percent) include Fe present at 48.00 to 88.00, Cr 0 to 32.00, Ni 0 to 16.00, Mn 0 to 21.00, B at 1.00 to 8.00 and Si at 1.00 to 14.00 atomic percent. The alloys maybe employed to provide a relatively wide range of strength and ductility depending upon whether or not such alloys form Class 1, 2 or 3 Steel as noted herein. The alloys may be used to form seamless tubular components by various methods, such as centrifugal casting.
After processing through centrifugal casting it is expected that the resulting pipe, tube, or tubular component may need to be post heat treated. The heat treatments can be one stage (such as 700 to 1200° C. for 5 minutes to 24 hours) or multiple stage (such as 850 to 1200° C. for 5 minutes to 24 hours, followed by a 200 to 1000° C. heat treatment for 5 minutes to 48 hours) depending on the property targets and the specific alloy response to the thermal exposure.
In the new alloys, melting occurs in one or multiple stages with initial melting from ˜1000° C. depending on alloy chemistry and final melting temperature might be up to ˜1500° C. Variations in melting behavior reflect a complex phase formation at chill surface processing of the alloys depending on their chemistry. The density of the alloys varies from 7.2 g/cm3 to 8.2 g/cm3. The mechanical characteristic values in the alloys from each Class will depend on alloy chemistry and processing/treatment condition. For Class 1 Steels, the ultimate tensile strength values may vary from 700 to 1500 MPa with tensile elongation from 5 to 40%. The yield stress is in a range from 400 to 1300 MPa. For Class 2 Steels, the ultimate tensile strength values may vary from 800 to 1800 MPa with tensile elongation from 5 to 40%. The yield stress is in a range from 400 to 1700 MPa. For Class 3 Steels, the ultimate tensile strength values may vary from 1000 to 2000 MPa with tensile elongation from 0.5 to 15%. The yield stress is in a range from 500 to 1800 MPa. Additional classes of steel are anticipated with possible yield strengths, tensile strengths, and elongation values outside of the limits listed above.
In the oil and gas industry there are four steel grades used for drill pipe and five common steel grades for casing and tubing whose properties are shown in Table 3 and Table 4. Properties for a selected Class 2 steel alloy, Alloy 82 (Table 2), from present application are included for comparison. Alloy 82 was cast in a copper die as a plate with thickness of 1.8 mm with subsequent thermal mechanical treatment.
Based on the data in Tables 3 and 4, Alloy 82 is close to being comparable to E75 drill pipe in yield strength yet is nearly twice as strong and would easily be comparable to J55 casing pipe and yet is over twice as strong. A stainless steel Alloy 82 is likely to have an additional advantage over both drill pipe and casing pipe grades in terms of higher corrosion resistance.
Because these tensile properties would be achievable in a pipe with a wall thickness of 2.1 mm (0.083 inch) there would be a significant weight savings for the different pipes as shown in Tables 5 and 6. Note that the largest outer diameter (OD) drill pipe from Alloy 82 would have lower in mass per unit length of the smallest diameter API spec drill pipe and a pipe from Alloy 82 would be nearly two and half to ten times lighter than any casing pipe in API spec. There would be a significant cost savings to drilling operations if the pipes used were substantially lighter. Additionally, much lighter pipe weight might be achieved by utilizing high strength Class 3 alloys or new classes of steel.
Raw feedstock material prepared in accordance with atomic ratios for Alloy 82 (Table 2) was heated in a furnace until melting. Upon homogenization, the molten liquid metal was transferred to a tundish that was moved to the mold location. The permanent mold was rotated continuously about its axis at a fixed speed as the molten metal was poured from the tundish. The molten metal was centrifugally thrown towards the inside mold wall where it solidified during cooling forming a pipe. The view of the cast pipe is shown in
The density of the pipe produced by centrifugal casting from Alloy 82 (Table 2) was measured using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3. The samples were cut by EDM from the pipe ring close to the outer diameter (OD) surface, close to the inner diameter (ID) surface and from the middle of the pipe wall. The density data are listed in Table 7 for each sample. As it can be seen, the density of the pipe is uniform through the volume.
Rockwell C hardness was measured on a cross sectional sample cut from the pipe produced by centrifugal casting from Alloy 82 (Table 2). Measurements were done by using a Newage Versitron manual Rockwell hardness tester and are listed in Table 8 and displayed as a function of distance in
Cross section plates with about 2 mm thickness were cut from the pipe produced by centrifugal casting from Alloy 82 (Table 2). The tensile specimens were cut from cross sectional plates using wire electrical discharge machining (EDM). The tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving; the load cell is attached to the top fixture. The tensile test results for the as cast pipe as a function of distance from the outer diameter (OD) surface are shown Table 9 with the tensile strength profile shown in
Tensile specimens cut from cross sectional plates of the pipe were heat treated from 1100° C. to 1200° C. for 1 hour and another set were subsequently heat treated at 1200° C. for 1.5 hours to see if the tensile properties improved. The as cast and heat treated tensile results are listed in Table 10 and shown in
SEM samples from the outer diameter (OD) region were cut from the pipe produced by centrifugal casting from Alloy 82 (Table 2). Samples were mechanically polished to a mirror finish and examined with a Zeiss Scanning Electron Microscope that has a maximum of 30kV accelerating voltage. The microstructure of the OD region is shown in
A prototype pipe was produced from Alloy 82 (Table 2) by hot extrusion of the billet consolidated from alloy powder by Hot Isostatic Pressing (HIP). Powder was produced by atomization process from the melt using industrial centrifugal atomizer. Raw feedstock material prepared in accordance with atomic ratios for Alloy 82 was inductively heated until melting in an inert atmosphere. Upon homogenization, the molten liquid metal was transferred to the reservoir of an industrial centrifugal atomizer. The molten liquid metal was poured continuously onto a rapidly rotating disc from which the liquid was aspirated into droplets that were rapidly cooled by an inert gas. The produced powder was collected in a chamber and processed through an air classifier that separated the powder at minus 180 μm (˜80 mesh). The alloy has good processability by powder atomization resulted in high quality powder with regular shape with yield of 60% under 180 μm in size.
The produced powder from Alloy 82 representing Class 2 steel was compacted by Hot Isostatic Pressing (HIP) method into a cylindrical billet using high pressure inert gas at elevated temperatures in order to consolidate encapsulated powders to give fully dense materials. The billet in a form of a cylinder with 10 inches in diameter and 13.5 inches in length was prepared for extrusion by boring a 4 in. dia. hole in the center. Additionally, a protective glass lubricant was applied to the surface to aid in extrusion and to prevent oxidation during soak in atmosphere at 1193° C. for 4 hr. The as-prepared billet is shown in
After soaking, the billet was removed from the oven and rolled across a woven glass cloth which adheres to the surface for additional lubrication. It was then moved to a lift which positions the billet between the hydraulic ram and the mandrel and extrusion die. During extrusion, the billet was forced through a die which reduces the outer diameter (OD) and holds the inner diameter (ID) constant by using a stationary 4 in. dia. mandrel. The material experienced an acceleration as it is forced through the die and was stretched longitudinally, reducing the cross sectional area and increasing the length. Both the die and the mandrel are preheated to 371° C. The extrusion process is illustrated in
The density of the pipe produced by hot extrusion from Alloy 82 (Table 2) was measured using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water. Experimental results have revealed that the accuracy of this technique is ±0.01 g/cm3. The density of the extruded pipe is 7.55 g/cm3 as an average of 5 measurements.
Rockwell C hardness was measured on a cross sectional sample cut from the pipe produced by hot extrusion of HIP consolidated billet from Alloy 82 powder. Measurements were done by using a Newage Versitron manual Rockwell hardness tester and are listed in Table 12. The hardness has similar values through cross section and ranges from 29.7 HRC to 33.1 HRC.
Cross section plates with about 2 mm thickness were cut from the pipe produced by hot extrusion from Alloy 82 (Table 2). The tensile specimens were cut from cross sectional plates using wire electrical discharge machining (EDM). The tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at room temperature in displacement control with the bottom fixture held ridged and the top fixture moving; the load cell is attached to the top fixture.
The pipe properties have been tested in the as-produced state, as well as after heat treatments at 900 for 1 hr and 1200° C. for 1.5 hours. Additionally, a study on the property values through the pipe thickness going from the internal diameter to the outer diameter has been undertaken. Tensile properties of the pipe material in different states are shown in
To evaluate a uniformity of the pipe properties through the volume, a series of tensile specimens was cut from a profile section of the extruded pipe, and then tested to determine if there is any variation in tensile properties through the pipe wall thickness from the inner diameter (ID) to the outer diameter (OD) of the pipe. All samples were heat treated at 900° C. for 1 hour. Despite some variation in property values, the properties are shown to be consistent through the thickness with most measurements above tensile strength of 1200 MPa and ductility more that 20% (
Charpy specimens were cut out from a pipe produced by hot extrusion of the billet compacted from Alloy 82 powder. Specimens were cut using wire electrical discharge machining (EDM) with two different orientations relative to the centerline axis of the pipe, which were identified as longitudinal and transverse as shown in
Samples for scanning electron microscopy (SEM) analysis were cut from as-extruded pipe at different locations through wall thickness. To ensure smooth surface, SEM samples were metallographically polished in stages down to 0.02 μm Grit. SEM was done using a Zeiss EVO-MA10 model with the maximum operating voltage of 30 kV manufactured by Carl Zeiss SMT Inc. As shown in
Example SEM backscattered electron micrographs of the as-extruded samples are shown in
To evaluate heat treatment effect on pipe microstructure, samples for SEM analysis were cut from heat treated pipe close to inner diameter, from the wall center and close to outer diameter of the pipe. Two heat treatments were applied to the pipe material: at 900° C. for 1 hour and at 1200° C. for 1.5 hrs. To ensure smooth surface, SEM samples were metallographically polished in stages down to 0.02 μm Grit. SEM was done using a Zeiss EVO-MA10 model with the maximum operating, voltage of 30 kV manufactured by Carl Zeiss SMT Inc.
To examine the structural details, high resolution transmission electron microscopy (TEM) was utilized. To prepare TEM specimens, samples were cut from the HIPed billet that was used to make the extruded pipe. Samples cut from the billet after high temperature heat treatment was also prepared for TEM studies. The samples were first ground and polished to a thickness of 50˜80 μm. Discs of 3 mm in diameter were then punched from these thin samples, and the final thinning was done by twin-jet electropolishing using a 30% HNO3 in methanol base. The prepared specimens were examined in a JEOL JEM-2100 HR Analytical Transmission Electron Microscope (TEM) operated at 200 kV.
TEM study suggests that the matrix is composed mostly of austenite phase (γ-Fe), and the grain size of the austenite is 1 to several microns in the as-HIPed billet. Finer grains produced through Static NanoPhase Refinement are sometimes seen in the as-HIPed billet, but their volume fraction is relatively small (
A potential level of properties and property combination range for alloys herein is shown in
Using commercial purity feedstock, charges of different masses were weighed out for Alloy 309, Alloy 327, Alloy 328, and Alloy 335 according to the atomic ratios provided in Table 2. The elemental constituents were weighed and charges were cast at 50 mm thickness using a Indutherm VTC 800 V Tilt Vacuum Caster. The feedstock was melted using RF induction and then poured into a water cooled copper die corresponding to cast process at commercial production. Subsequent hot rolling was applied using a Fenn Model 061 Rolling Mill and a Lucifer 7-R24 Atmosphere Controlled Box Furnace. The samples were hot rolled to ˜96% reduction in thickness via several rolling passes following a 40 minute soak at 50° C. below each alloy's solidus temperature. Alloy 309 and Alloy 327 were heat treated at 850° C. for 6 hr. Alloy 335 was heat treated at 850° C. for 6 hr. Alloy 328 was tested in hot rolled condition.
It should also be noted that the alloys formed herein may be employed in other configurations, such as in the form of sheet. Such sheet may have a thickness from 0.3 mm to 150 mm and from 100 mm to 5000 mm in width. The alloys herein in either forms as Class 2 or Class 3 Steel therefore have a variety of applications. These include but are not limited to structural components in vehicles, including but not limited to parts and components in the vehicular frame, front end structures, floor panels, body side interior, body side outer, rear structures, as well as roof and side rails. While not all encompassing, specific parts and components would include B-pillar major reinforcement, B-pillar belt reinforcement, front rails, rear rails, front roof header, rear roof header, A-pillar, roof rail, C-pillar, roof panel inners, and roof bow. The Class 2 and/or Class 3 steel will therefore be particular useful in optimizing crash worthiness management in vehicular design and allow for optimization of key energy management zones, including engine compartment, passenger and/or trunk regions where the strength and ductility of the disclosed steels will be particular advantageous.
The alloys herein may also provide for use in additional non-vehicular applications, such as for drilling applications, which therefore may include use as a drill collars (a component that provides weight on a bit for drilling), drill pipe (hollow wall pipe used on drilling rigs to facilitate drilling), pipe casing, tool joints (i.e. the threaded ends of drill pipe) and wellheads (i.e. the component of a surface or an oil or gas well that provides the structural and pressure-containing interface for drilling and production equipment) including but not limited to ultra-deep and ultra-deep water and extended reach (ERD) well exploration. The alloys herein may also be used for a compressed gas storage tank and liquefied natural gas canisters.
This application claims the benefit of U.S. Provisional Application Ser. No. 61/750,606 filed Jan. 9, 2013.
| Number | Date | Country | |
|---|---|---|---|
| 61750606 | Jan 2013 | US |
