US20160114462A1

US20160114462A1 - Apparatus and method for subsurface structural modification of materials at reduced temperatures

Info

Publication number: US20160114462A1
Application number: US14/923,878
Authority: US
Inventors: Laszlo J. Kecskes; Micah J. Gallagher; Anthony J. Roberts; Heather A. Murdoch; Kristopher A. Darling
Original assignee: United States, AS REPRESENTED BY SEC; US Army Research Laboratory; US Department of Army
Current assignee: BOWHEAD SCIENCE AND TECHNOLOGY LLC; United States, AS REPRESENTED BY SEC; US Army Research Laboratory; US Department of Army
Priority date: 2014-10-27
Filing date: 2015-10-27
Publication date: 2016-04-28

Abstract

Nanostructured or ultra-fine grained metallic systems according to embodiments of the invention may be formed of: pure Cu, pure Fe, or pure Ti, with grain sizes of less than 140 nm, 348 nm, or 59 nm, respectively. The metallic systems demonstrate a monotonically increasing grain size dependence from a minimum value attained at the surface; and a converse relation of microhardness, decreasing from 160 kg/mm², 265 kg/mm², or 320 kg/mm², respectively. The grain refinement process at cryogenic conditions relies on the suppression of room temperature dislocation-mediated deformation mechanisms which facilitate grain restructuring, relaxation, and reorientation. At the cryogenic conditions, alternative mechanism for grain refinement, such as shear localization or dynamic recrystallization may be more dominant. Processes for refining the grain size of these metallic systems may include: subjecting metal plates to a high-energy milling process using a high-energy milling device to impart high impact energies to its surface. Due to the high-efficiency of this attrition process, these metallic systems are ideal candidates for improved corrosion and wear resistance.

Description

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims priority to U.S. Provisional Application No. 62/068,994, titled “Apparatus and Method for Subsurface Structural Modification of Materials at Cryogenic Temperatures,” filed on Oct. 27, 2014 which is hereby incorporated by reference herein including all attachments and papers filed with U.S. Provisional Application No. 62/068,994.

GOVERNMENT INTEREST

The invention described herein may be manufactured, used, and licensed by or for the United States Government without the payment of royalties thereon.

BACKGROUND

1. Field of the Invention
The present disclosure relates to nanostructured surface coatings and subsurface modifications, having uniform to gradient microstructures, with single or multiple elemental constituents, and methods of making the same.
2. Background of the Invention
Nanostructured and nanocrystalline materials are well known to have novel and/or enhanced properties compared to their coarse-grained counterparts. Specifically, nanostructured metals may have enhanced strength, hardness, and/or improved workability. It has been shown that these nanostructures, when formed as metallic surface layers, could result in a number of significant improvements and new applications. For example, failure in metallic materials, such as fatigue fracture, fretting fatigue, or wear and corrosion often initiate and propagate from the surface of the material. As such, optimization by nanostructuring of the surface layer through surface modification would limit such detrimental behavior and, in turn, not only enhance the surface, but also the global or bulk properties of the material.
However, the abrupt transition in physical characteristics from such a surface layer to the bulk would likely result in a correspondingly different response to physical phenomena as well. Therefore, it is hypothesized that creation of a gradual transition region between the unaltered bulk and this surface coating, having a range in properties, would be more effective in mitigating such phenomena.
There are a number of techniques to generate nanostructured surfaces by using traditional coating techniques, e.g., plasma vapor deposition or chemical vapor deposition.
In the recent decade, similar to the well established method of shot peening, a novel Surface Mechanical Attrition Treatment (SMAT) technology has been successfully applied to many metal-based material systems such as those based on Ni, Fe, Mg, Al, or Ti.
Surface mechanical attrition treatments have received much attention due to their ability to enhance physical properties, such as yield strength, hardness, wear, and fatigue resistance in structural parts. Fundamentally, such treatments are carried out by repeated impacts of the metal surface with some impingement media which normally occurs at high strain rates inside a closed or partially open container. Partially inelastic collisions during processing transfer kinetic energy into the substrate in the form of deformation strain energy and heat. In the initial stages of SMATing it is expected that, at the surface, localized deformation occurs through compressive stresses and shear banding followed by heavy deformation, wherein the total strain energy increases due to the creation of excess dislocations, stacking faults, and other subgrain structures leading to an eventual attenuation and reduction of the grain size. Finally, an equilibrium grain size is reached, the specific length scale of which is determined by a balance between the flux of induced defects, the evolution of the generated defect cell structures, and their relaxation, reorganization, and restructuring. As relaxation/restructuring is thermally activated, the local as well as global temperature rise during processing plays a significant role in determining the resultant microstructure during SMAT processing.
The prior art teaches the conventional application of SMAT, wherein, plastic deformation, at high strains and strain rates, is imposed on the surface layers of bulk samples. Typically, spherical steel (or ceramic) balls of a few millimeters in diameter with smooth surfaces are placed in the bottom of a chamber that is sonically vibrated (usually with a frequency of 50 Hz˜20 kHz). When the balls are resonated, the sample surface is impacted by a large number of flying balls over a short period of time. Each impact of a flying ball (with a velocity of 1˜20 m/s) induces plastic deformation at a high strain rate in the surface layer. Repeated multidirectional impacts result in repeated plastic deformation in the top surface layer and the creation of bulk structural defects (dislocations, subgrain and stacking fault formation, etc.) followed by a process of grain restructuring, relaxation, and reorientation, which, in turn, induces progressive grain refinement down to the nanometer regime. What differentiates this form of SMAT from shot peening is that in the latter case, the balls are directed onto the surface using a nozzle. As such, the kinetic energy and directionality of each ball is roughly the same as another ejected from the nozzle, rendering no major variation in the total quanta of deformation energy being imparted to the surface.
Surface mechanical attrition treatment offers some unique advantages over coating and deposition methods. For one, there is no change in chemical composition from the surface to the bulk or matrix. Another advantage is the relative ease in the ability to create a continuous, gradient variation from the top surface, which is nanocrystalline, less than 100 nanometers, to the sub-surface, which is ultra-fine grained, less than one micrometer, to the bulk, which is coarse-grained, greater than one micrometer. With SMAT, bonding between the surface and matrix is not an issue.
The observed surface attrition process presents similar to that of mechanical alloying/milling during active communication where the imparted kinetic energy depends on the frequency and the amplitude of the specific mill used; and, like in mechanical milling, a similar analogous temperature rise (50-150° C.) in the substrate is expected to occur for higher energy processes. This is primarily due to the deformation induced heat generated within the substrate, ball to ball, ball to wall collisions, and frictional heating effects during the repeated attrition process. While generally this temperature rise is considered small, it can have a noticeable effect on certain metals and alloys (e.g., high purity, face centered cubic [FCC] nanocrystalline metals which have been shown to undergo grain growth at room temperature).
Here we describe the procedures and advantageous effects of cryogenic SMAT processing at sub-ambient temperatures, preferably below the ductile-brittle transition temperature of the material being treated, more preferably, below −50° C., most preferably, at liquid nitrogen temperature for commercially available oxygen free high conductivity copper (Cu) (OFHC), pure iron (Fe), and pure titanium (Ti), each representative of a different crystallographic system. In the spirit of this invention, these exemplary systems are used as illustrations for what is possible with our concept. However, it is noted and emphasized that the applicability of the processes described herein have greater utility to a much broader range of material systems, including, but not limited to, other pure metals and alloys with commercialization potential. In turn, to demonstrate the unique, non-obvious, and salient features of this invention, we will differentiate the low temperature structural evolution of these metals from that occurring during room temperature processing.
Note, while SMAT processes have been applied to a number of systems including, essentially pure metals such as Cu, Ni, Ti, Fe, and alloys such as stainless steels, currently, there is no or very limited data on decoupling the thermomechanical effect during the processing of such materials, especially at low temperatures where dynamic recovery and recrystallization are dramatically suppressed.
This invention utilizes equipment, developed for high energy cryogenic mechanical alloying/milling, which presents a unique, non-obvious opportunity when applied to low temperature or cryogenic SMATing.

BRIEF SUMMARY OF THE INVENTION

Various unary, binary, and higher order metallic systems, and methods of making the same, are presented herein according to embodiments of the invention.
According to various embodiments, these metallic systems may include: pure copper (Cu), pure iron (Fe), or pure titanium (Ti), each representative of a specific crystallographic class, face centered cubic (FCC), body centered cubic (BCC), and hexagonal close packed (HCP), respectively. The untreated coarse-grained metal, upon being subjected to cryogenic SMAT, will possess a gradient microstructure from the surface to the bulk. That is, the as-processed metal will consist of a nano- to submicrometer length scale grain structure at its surface, monotonically reaching, some distance away below the surface, the initial, untreated grain size of the bulk. The ability to control the depth below the surface and the effectiveness of this grain size reduction is one the novel aspects of this invention.
According to various embodiments, these metallic systems may also include: multiple elements, wherein upon SMAT, the resultant component elements may be in a discrete or disordered or alternatively ordered, layered, in a stratified arrangement, consisting of a single to multiple layers, each with its own distinct features, displaying various levels of intermixing between the initial constituents. Alternatively, adjustment of the conditions may generate a cellular structure on a multiscale level, with varying dimensions of the cell size.
In the various embodiments of the metallic systems, the initial metal may have an untreated grain size greater than 20 to 50 μm. The various exemplary metal systems, subjected to room temperature SMAT, may have grain sizes as fine as 150 nm up to 360 nm. However, the grain size reduction, only after one hour treatment, across the three exemplary metallic systems, is considerably more effective at cryogenic temperatures: 60 nm up to 140 nm, or 45% to 62%, respectively.
For the purposes of this invention, various ASTM measurement standards or their accepted equivalents were used. For example, for the determination of microhardness, ASTM E384-11e1, Standard Test Method for Knoop and Vickers Hardness of Materials, ASTM International, West Conshohocken, Pa., 2011 was used. Likewise, for the determination of grain size, the linear intercept method, quite common in microstructural analysis practices was used as well as the ASTM E112-13, Standard Test Methods for Determining Average Grain Size, ASTM International, West Conshohocken, Pa., 2013 was used as reference.
Furthermore, in these various embodiments, the depth of nanostructuring, as measured from the surface, can vary from 250 to 1000 μm. With the use of cryogenic temperatures, the depth changes, by as much as a factor of two, as determined by the crystallographic class of and the ease of deformation within the material. As such, for the BCC and HCP cases, the depth of the affected region is significantly narrower, whereas, for the FCC case, it is wider. While it is preferred that a low temperature is used, more preferably the lowest attainable with the cooling system, various embodiments may necessitate altering and increasing the processing temperature to prevent the processed material from prematurely fracturing or failing, due to it being too brittle or below its ductile-brittle transition temperature.
These embodiments thus provide a new class of nanostructured and nanocrystalline metals and alloys which have significantly improved properties compared to their coarse grained counterparts. Moreover, the exemplary copper (Cu) metallic system subjected to room temperature SMAT may have a Vickers microhardness (HV) of about 110 kg/mm²at its surface, while an equivalent sample, processed at a cryogenic temperature, may have an HV of about 155 kg/mm²; a significant increase. Similarly, the exemplary iron (Fe) metallic system subjected to room temperature SMAT may have an HV of about 220 to 250 kg/mm²; for iron (Fe) the resultant hardness difference between room and cryogenic temperature processed samples is not notable. Lastly, the exemplary titanium (Ti) metallic system, processed with room temperature SMAT may have an HV of about 280 kg/mm²; whereas, the cryogenically treated sample may have an HV of about 320 kg/mm². Again, in the spirit of this invention, these aforementioned microhardness values are illustrations of typical properties for a specific processing condition, one hour of SMAT. It is conceivable that longer SMAT times will lead to further hardness increases, especially at the surface. It is also believed that shorter SMAT times, i.e. impact times, may be sufficient to achieve the desirable improvement in final properties. Thus, in certain embodiments, the processing time, i.e. the total amount of time that the metal parts or parts are impacted with metal fragments (t_i) at reduced temperature (T_r) is at least about 5 minutes, more preferably at least about 10 minutes, still more preferably at least about 20 minutes, still even more preferably at least about 40 minutes and in some instances at least about 1 hour or 60 minutes.
According to further embodiments, the SMAT process for forming the exemplary nanostructured Cu, Fe, or Ti metallic systems, comprised of a gradient microstructure from nano- to micro- to macroscales, imparts these metals with a preferred grain orientation or texture. For each system, the texture evolves, as determined by the ease of deformation and grain reorientation when subjected to high strain and high stress conditions.
During SMAT, using a high-energy milling device may utilize a mixing vial to contain the metallic plate in addition to or replacing the vial's lid and, in some circumstances, may utilize a plurality of milling balls for inclusion within the mixing vial. In other embodiments, especially for those when surface alloying is desired, in addition to the plurality of milling balls, the vial may contain the powder(s) to be alloyed therein. In certain circumstances, the vial and steel milling media may be precoated with a specific metal to preferentially augment or reduce the presence of a specific element or species. The ball-to-powder mass ratio utilized by the high-energy milling device may be 1:1, more specifically, 10:1, or more. Furthermore, the milling balls may be comprised of high strength steel or ceramic. During the high-energy milling process, the vial and its contents may be cooled to sub-ambient temperatures, more preferably, a cryogenic temperature. This may be accomplished by cooling the milling device with any type of cryogenic liquid, more specifically, liquid nitrogen. Alternatively, the high-energy milling process may be performed at ambient or room temperature. The high-energy milling process may be further improved using an additive or a surfactant. In some instances, the vial and its contents may be continuously or semi-continuously cooled during the high-energy milling process. In further embodiments, at the conclusion of the milling process, the metallic powder may be subjected to annealing by exposing it to elevated temperature in the range of about 300 to 800° C.
These embodiments thus provide a methodology for forming a new class of nanostructured and nanocrystalline metals or composites which possess enhanced mechanical properties with a gradient grain size microstructure.
These and other, further embodiments of the invention are described in more detail, below.

BRIEF DESCRIPTION OF THE DRAWINGS

So that the manner in which the above recited features of the present invention can be understood in detail, a more particular description of the invention, briefly summarized above, may be had by reference to embodiments, some of which are illustrated in the appended drawings. It is to be noted, however, that the appended drawings illustrate only typical embodiments of this invention and are therefore not to be considered limiting of its scope, for the invention may admit to other equally effective embodiments, including less effective but also less expensive embodiments which for some applications may be preferred when funds are limited. These embodiments are intended to be included within the following description and protected by the accompanying claims.

FIG. 1 shows scanning electron micrographs of the gradient nature of the subsurface grain size reduction in pure copper (Cu), comparing samples processed at ambient and cryogenic temperatures, respectively.

FIG. 2 shows the x-ray diffraction patterns of the as-received, ambient temperature processed, and cryogenically processed copper (Cu) samples, respectively.

FIG. 3 shows scanning electron micrographs of the gradient nature of the subsurface grain size reduction in pure iron (Fe), comparing samples processed at ambient and cryogenic temperatures, respectively.

FIG. 4 shows the x-ray diffraction patterns of the as-received, ambient temperature processed, and cryogenically processed iron (Fe) samples, respectively.

FIG. 5 shows scanning electron micrographs of the gradient nature of the subsurface grain size reduction in pure titanium (Ti), comparing samples processed at ambient and cryogenic temperatures, respectively.

FIG. 6 shows a high magnification scanning electron micrographs for copper (Cu), iron (Fe), and titanium (Ti) metal types revealing the dramatic decrease in grain size with the use of cryogenic processing conditions versus those at ambient conditions.

FIG. 7 shows the x-ray diffraction patterns of the as-received, ambient temperature processed, and cryogenically processed titanium (Ti) samples, respectively.

FIG. 8 shows the Vickers microhardness values of the copper (Cu), iron (Fe), and titanium (Ti) samples, as a function of depth away from the surface for both the ambient and cryogenic temperature processed samples.

DETAILED DESCRIPTION OF INVENTION

Unlike that in the prior art, we utilize for the first time a novel low temperature or cryogenic SMAT process on commercially available, oxygen-free high-conductivity (OFHC) copper (Cu), pure iron (Fe), and pure titanium (Ti) and show the advantages in terms of microstructure refinement as compared to conventional ambient or room temperature (RT) SMAT processing. While SMAT processes in a conventional configuration have been applied to a number of metallic systems, including Cu, Ni, Ti, Fe, and stainless steels, there currently is very limited data demonstrating decoupling of thermomechanical effects during the processing, especially at cryogenic temperatures (e.g., below −150° C.) where dynamic recovery and recrystallization are dramatically suppressed. By utilizing equipment developed for high-energy cryogenic mechanical alloying or milling, which is not typically used, it is possible to show the substantial microstructure refinement of the cryogenic SMAT process in OFHC Cu, pure Fe, or pure Ti over what is attainable at RT.
Thus, in certain preferred embodiments the present invention provides a method of modifying the surface of a metal part that includes: providing at least one metal part that is formed from a first metal composition; providing a plurality of metal fragments that are formed from a second metal composition wherein said metal fragments have a size that is significantly smaller than the size of the at least one metal part; reducing the temperature of the at least one metal part and the plurality of metal fragments; impacting the at least one metal part with the plurality of fragments at a reduced temperature (T_r), for an impact processing time (t_i); wherein, the at least one metal part is subjected to bombardment by the plurality of metal fragments at a reduced temperature (T_r), resulting in the as-received grain size of the at least one metal part to be reduced, by several orders of magnitude, and possessing a gradient structure from the impact surface of the at least one metal part into the interior of the bulk of the at least one metal part.
In the exemplary embodiments provided below, the metal fragments were steel balls or milling balls. However, ceramic milling balls may be used or some other material exhibiting properties similar in nature. Suggested steel milling balls will have a diameter of at least about 0.015 mm and more preferably at least about 1.5 mm and may have a diameter as large as about 2.5 cm. In certain embodiments, the milling balls are formed from tungsten. Thus, the mass of an individual milling ball may range from about 1.01 gram to about 160 grams. Desirably, the properties of the milling balls, including but not limited to hardness, should be similar to or greater than the properties of the material being milled or SMATed. However, in certain embodiments the milling balls may be abraded during the process to produce an alloy or composite structure at the surface of the part or material that is being milled or SMATed. In certain other embodiments a powder of is included in the milling container to produce an alloyed surface or to introduce a second phase at the SMATed surface.
In the exemplary embodiments provided below, the at least one metal part included a plurality of metal plates. However, in commercial applications it is suggested that the at least one metal part is a finished part or an almost finished part, for example a head of a golf club, a gear or gear tooth or a part of a gun or an engine component prior to assembly or packaging. The composition of the metal part(s) may be same or different from the composition of the metal fragments. In the exemplary embodiments provided below, the composition of the metal parts, e.g. metal plates, and the composition of the fragments were different. However, in some applications it is may be desirable that the composition of the metal part(s) may be same as the composition of the metal fragments or substantially the same as the composition of the metal fragments. For example, both the part(s) and the impact fragments may be formed from high-conductivity (OFHC) copper. In certain embodiments the surface at least one part that is being SMATed at reduced temperature is cut away and then used as a sub part or component. Additionally, the SMATed surface may be further processed. For example, the high density of grain boundaries in a nanocrystallline microstructure also increases the available diffusion pathways at the surface for secondary processing such as carburization or nitriding, enabling decreases in process temperature of possibly improved species absorption.
Samples were prepared by cutting commercially available OFHC Cu, pure Fe, or pure Ti into disks 6.35 cm in diameter and 0.6 cm thick. These samples were then polished to a mirror finish. Nominally supplied high energy SPEX® SamplePrep Corporation mill vial lids were then replaced by the samples and sealed in a high-purity argon glovebox. In each case, the hardened steel vials were loaded with 17 5/16″ diameter and 16¼″ diameter 440 C stainless steel ball bearings to constitute a total mass of 50 g. For the Cu and Ti samples, the vial was precoated with a thin layer of Cu or Ti, respectively. This was to limit potential contamination of the as SMATed surface from the Fe constituent of the milling vial. The precoating was accomplished by placing approximately 0.5 g of each respective powder into the vial and operating the SPEX mill for 10 min at room temperature. The actual SMAT samples were either processed at cryogenic or RT for a period of 1 hour. For the purposes of the invention, low or cryogenic temperatures are defined as temperatures well below ambient room temperature conditions, preferably below about −50° C. (223K), not greater than about −100° C. (173K), more preferably not greater than about −150° C. (123K), and most preferably not greater than about −196° C. (77K). RT milling was accomplished by loading the vials into a commercially available single vial SPEX SamplePrep Model 8000M mill, while cryogenic milling was performed in a modified SPEX SamplePrep Model 8000M mill. The modified mill was equipped with a Teflon sleeve into which the sealed steel vials could be inserted. The Teflon sleeve was fitted to allow the inflow and outflow of liquid nitrogen at a temperature of −196° C. (77K) to envelope the outside of the steel vial. After the cryogenic and RT SMAT processing, the samples were mounted and polished for microstructure analysis.
The microstructure was analyzed using optical microscopy, X-ray diffraction for grain size and strain analysis, and scanning electron microscopy (SEM) for chemical, grain size and orientation analysis. Electron backscatter imaging in the SEM was used to analyze the grain size. Optical imaging of etched samples was used to assess the specific grain size and defect morphology. A dual-beam FEI Nano600 FIB was used to prepare samples for electron backscatter diffraction (EBSD) imaging and produce ion-channeling contrast images to highlight the active deformation mechanisms during cryogenic SMAT processing.
FIGS. 1a and 1b are SEM images of the etched sample cross-sections of two copper (Cu) samples prepared by RT and cryogenic SMAT processing for 1 hour, respectively. For the purposes of the invention, the duration of processing is defined as times being at least 5 minutes, preferably at least 10 minutes, more preferably at least 30 minutes, and most preferably at least 60 minutes. Note, however, the duration of processing is a strong function of the ductile-brittle temperature or malleability of the metal being treated, its crystallographic class, and the temperature dependence of its malleability. As such, for the purposes of the invention, there is no upper bound in processing times, except that determined by the failure of the metal plate due to void formation, fracture, spallation, and subsequent fragmentation. As such, under certain conditions, processing times as long as two to three hours may be feasible. The exemplary and representative micrographs illustrate the change in grain structure as a function of depth from the SMAT surface, which is analogous to how the microstructure evolves as a function of time for the SMAT process. The SMAT surface had up to 20 μm non-continuous layer with some Fe contamination from the use of stainless steel ball bearings; this is present in both RT and cryogenic samples.
There is a noticeable transition from nanostructured or ultrafine-grained grain structure to a region of banded structures to the bulk microstructure within the first 500 μm. The back surface of the SMAT specimens (images to the left, denoted by a square) is 0.6 cm from the SMAT surface. The microstructure near the back surface represents a pristine microstructure unaffected by the SMAT process; the average equiaxed grain size in this region is 150 μm in diameter with some twinning present. The average grain size for the cryogenic sample (140 nm mean diameter) is approximately 60% smaller than the grain size achieved through RT processing (355 nm mean diameter).
The cryogenic SMAT sample maintains a constant finer grain size to a much deeper penetration depth. That is, in contrast to the cryogenic sample, the RT processed sample appears to have a continuous monotonic increase in grain size as a function of depth into the plate. Furthermore, the overall region of grain refinement, defined here as the volume containing grains smaller than 10 μm in diameter, is found to vary with processing conditions. This region is approximately 150 μm deep for the cryogenically processed sample and 300 μm deep for the RT processed sample, respectively.
As shown in FIG. 2, X-ray diffraction patterns were collected from the as-received OFHC copper (Cu) plate, RT, and cryogenic SMAT samples. Upon inspection for the fundamental Bragg crystallographic reflections, the full widths at half-maximum increased and the amplitude decreased in the following order: as-received, 1 hour RT condition and 1 hour cryogenic condition. These trends are consistent with a decrease in grain size and/or an increase in local strain with processing conditions. The Scherrer estimates of grain size indicate that the average grain size for the cryogenic SMAT samples is lower than that calculated for the RT SMAT samples. However, the grain size estimates using the Scherrer formula were significantly lower than the previously measured grain size using SEM micrographs.
The texture change during SMAT can also be measured by the change in the relative intensity of the fundamental crystallographic reflections of Cu. For the cryogenic SMAT sample, the relative intensities of the fundamental reflections of pure Cu, as given by the JCPDS index, are: (111) 100%, (200) 46%, (220) 20%, (311) 17% and (222) 5%. For the unprocessed plate, the second and third reflections have higher percent relative intensities: (200) 58% and (220) 63%. This is in contrast to the texture generated during the SMAT process. Hence, in general, SMAT induces a texture wherein the (111) orientation is favored relative to the other reflections. Furthermore, cryogenic SMAT processing enhances this type of texturing over RT SMAT processing. For the cryogenically processed sample, the ratio of intensities of the (200) and (111) reflections increases to a slightly higher value than that with RT processing.
FIGS. 3a and b show the equivalent SEM images of the etched iron (Fe) sample cross-sections processed for 1 hour, respectively. In these images the contrast in grain morphology between the RT and cryogenic processed samples is rather striking. For iron (Fe) processed at RT, the deformed and reduced grains show a significant texturing effect, wherein the grains appear in layers and are stratified perpendicular to the plate normal. Whereas, the iron (Fe) processed at a cryogenic temperature, the as-deformed region is less stratified and more equiaxed. Comparing the affected depths between conditions, the RT processed specimen is significantly deeper than that processed at cryogenic temperature. For the former, the highly refined grain size region is about 200 μm deep; for the latter, the corresponding region is only about 50 μm. An intermediate region between the unaltered bulk is also wider for the sample processed at RT.
Unlike the FCC copper (Cu) samples which showed a dramatic increase in the (111) crystallographic orientation relative to all of the others, the BCC iron (Fe) samples show a considerably lesser effect. As shown in FIG. 4, are the X-ray diffraction patterns of the as-received iron (Fe) plate, RT, and cryogenic SMAT samples. Inspection of the fundamental Bragg reflections show that while the full widths at half-maximum did significantly change, the amplitudes did not decrease between the conditions presented. This effect is directly related to peak broadening due to grain size reduction, however, without, a significant change in crystallographic orientations. That is, aside from a reduction in grain size and pancaking of the grains, SMAT, in either RT or cryogenic conditions, does not introduce grain reorientation.
As was seen for both copper (Cu) and iron (Fe), there is a corresponding decrease in the size of region of nano to ultrafine-scale grains for HCP titanium (Ti) as well. The grains appear to be mostly equiaxed; absent is the layering or stratification of the grains. In FIGS. 5a and 5 b, SEM images of the etched sample cross-sections of two titanium (Ti) samples prepared by RT and cryogenic SMAT processing for 1 hour, respectively, show that submicrometer grains persist to a depth of about 75 μm in the RT processed sample, whereas, this depth is only about 40 μm for the cryogenic sample. Again, the transition region from the nano-scale to macroscale grains is deeper for the cryogenic condition. However, overall, the depths of the affected regions are about the same; 450 μm. This effect is similar to the case of copper (Cu), but different from that of iron (Fe).
X-ray diffraction patterns of the as-received titanium (Ti) plate, RT, and cryogenic SMAT samples are shown in FIG. 6. The full widths at half-maximum of the Bragg peaks illustrate significant line broadening for both SMAT conditions. In fact, the two primary reflections of (002) and (101), partially and completely overlap for the RT and cryogenic SMAT conditions, respectively. Note, with the exception of the growth of the (101) peak, the there is little change in the relative peak heights of the other peaks between the processing conditions. Given the fact that HCP titanium (Ti) has a limited number of operative deformation slip systems it is most likely, that recrystallization is the primary grain refinement mechanism in this material. This is consistent with the observed grain morphology and lack of texturing as indicated by the changes in relative peak heights.
For the purposes of the invention and the preferred embodiments described herein, it is important to realize that, whereas, the one hour processing time may have been closer to optimum conditions for FCC copper (Cu), however, this processing time may not have been the case for latter, namely the BCC iron (Fe) and HCP titanium (Ti) embodiments. That is, it is implied from a comparison of results that the resultant texturing in the latter systems have not yet fully evolved. In other words, different crystallographic systems will develop differently due to their intrinsic properties and underlying deformation mechanisms. Thus, it is likely that longer processing times would have resulted in the evolution of a stronger texture in the other exemplary metals.
Another factor in these embodiments is the evolution of steady state equilibrium conditions during SMAT processing, wherein, the heat generation due to deformation is offset by active cooling. That is, in all likelihood, the processing temperature, while fixed at −196° C. (77K), primarily, for convenience, delivery, and availability of liquid nitrogen will have a significant effect on the effectiveness of the grain size refinement process. Thus, for the purposes of this invention, it is hypothesized that it is highly likely that at lower or higher temperatures, corresponding to alternate equilibrium conditions, would result in a different more favorable outcome, i.e., potentially finer grain size reduction for the latter metal systems.
FIG. 7 reveals higher magnification SEM images of regions located at greater depths below the SMAT surface. There are major microstructure differences with respect to the processing conditions. First, the cryogenically processed samples maintain a near-equiaxed grain morphology, whereas the RT processed samples have larger regions where the grain morphology is distorted from its initial equiaxed shape. Second, a high density of etch pits, mostly likely associated with dislocations intersecting the polished surface, are present in the RT sample. When metals undergo severe plastic deformation at or near RT, dislocation slip and deformation twinning are the principal modes of deformation. While both FCC and BCC metals have adequate numbers of operational slip systems, there is a considerable limitation of such slip systems in HCP metals. Regardless, dislocation tangles and their specific subgrain structures (dislocation cells, walls, geometrically necessary boundaries and incidental dislocation boundaries) are generated and equilibrated by thermally activated processes. As such, these are the means for achieving grain refinement during SMAT processing.
The cryogenic SMAT samples, however, show very different microstructure evolution to that of the RT SMAT samples. At lower temperatures, the dislocation-based processes are suppressed, as indicated by the lack of etch pits in the cryogenically processed samples. Moreover, in contrast to the RT SMAT samples, a large number of banded structures were observed within the large grains of the cryogenic SMAT samples. In many cases, these banded structures initiate from the grain boundaries and either terminate at the inside wall of the same grain or intersect the opposite side boundary. The spacing between many of the bands is less than 10 μm. In general, these bands contain small ultrafine equiaxed grain structures or structures that are elongated and parallel to each other across the width of the band. These observations are consistent with observations of twin/matrix bundles, bamboo nanograins, and shear bands that evolve as a function of strain during dynamic plastic deformation under liquid nitrogen temperatures. These observations indicate a shift in the dominant deformation mechanism during SMAT from dislocation-mediated behavior at room temperature to twinning/shear band-mediated at cryogenic temperatures.
A time evolution of the microstructure can be constructed for the cryogenic SMAT process. First, due to the high-strain-rate impacts at cryogenic temperatures, twin/matrix bundles and/or shear bands originate at the surface or at internal stress concentration sites (grain boundaries, triple junctions, etc.). The repeated impacts generate overlapping twin/matrix bundles and/or shear bands, which effectively refine the large micrometer-sized grains. Additional grain refinement may occur within any shear band present due to dynamic recrystallization. As the internal stress continues to accumulate, the microstructure evolves by producing ultrafine grains with long aspect ratios. The accumulated shear strain leads to fragmentation and rotation of these elongated grains, resulting in an equiaxed nanocrystalline microstructure.
FIG. 8 displays the results of microhardness measurements taken from the three metal specimen types both after being subjected to RT or cryogenic SMAT. The three embodiment exhibit a monotonic fall off in microhardness that strongly depends on the local grain size at the relative depth measured from the surface. Whereas, the RT induced hardness change is gradual from the surface to the interior, for samples processed at cryogenic temperatures, this hardness change is more rapid, presumably due to the shallower nature of the structurally modified region. The most significant change occurs for pure copper (Cu), then pure titanium (Ti), and lastly pure iron (Fe). The figure clearly illustrates the variations between the three embodiments. The advantage for the use of cryogenic conditions is most beneficial in the case of the softer copper (Cu), however, if a sharper hardness profile and grain size restructuring is desired, this could be achieved by altering the SMAT conditions, with lengthening the duration and potentially lowering the temperature further. The embodiment using titanium (Ti) illustrates that quite well. Note, in the case of pure iron (Fe), there is little to no benefit to the use of cryogenic processing conditions, below the surface. However, this does not preclude the use and the resultant benefits that derive from lower temperatures or longer SMAT times.
Supporting and/or additional details may be found in a journal article titled “Enhancing grain refinement in polycrystalline materials using surface mechanical attrition treatment at cryogenic temperatures” Scripta Materialia 69 (2013) 461-464 by Dr. Kristopher A. Darling et al. which is hereby incorporated by reference herein.
The foregoing description, for the purposes of providing an explanation, has been described with reference to specific embodiments. However, the illustrative discussions above are not intended to be exhaustive or to limit the invention to the precise forms and conditions disclosed. Many modifications and variations are possible in view of the above teachings. The embodiments were chosen and described in order to best explain the operating principles of the present disclosure and its practical applications to thereby enable others skilled in the art to best utilize the invention and its various embodiments with various modifications as may be suited to the particular use contemplated.
While the foregoing is directed to embodiments of the present invention, other and further embodiments of the invention may be devised without departing from the basic scope thereof, and the scope thereof is determined by the claims that follow.

Claims

1. A surface mechanical attrition treatment process for a metal part comprising:

providing at least one metal part that is formed from a first metal composition;

providing a plurality of metal fragments that are formed from a second metal composition wherein said metal fragments have a size that is significantly smaller than the size of the at least one metal part;

reducing the temperature of the at least one metal part and the plurality of metal fragments;

impacting the at least one metal part with the plurality of fragments at a reduced temperature (T_r), for an impact processing time (t_i);

wherein, the at least one metal part is subjected to bombardment by the plurality of metal fragments at a reduced temperature (T_r), resulting in the as-received grain size of the at least one metal part to be reduced, by several orders of magnitude, and possessing a gradient structure from the impact surface of the at least one metal part into the interior of the bulk of the at least one metal part.

2. The surface mechanical attrition treatment process of claim 1, wherein the reduced temperature (T_r) is not greater than about −50° C.

3. The surface mechanical attrition treatment process of claim 1, wherein the reduced temperature (T_r) is not greater than about −100° C.

4. The surface mechanical attrition treatment process of claim 1, wherein the reduced temperature (T_r) is not greater than about −150° C.

5. The surface mechanical attrition treatment process of claim 1, wherein the reduced temperature (T_r) is not greater than about −196° C.

6. The surface mechanical attrition treatment process of claim 1, wherein the refined grain size of the metallic system at its surface is about 1000 nm or less.

7. The surface mechanical attrition treatment process of claim 3, wherein the room temperature refined grain size of the metallic system at its surface is about 355 nm or less.

8. The surface mechanical attrition treatment process of claim 3, wherein the cryogenic refined grain size of the metallic system at its surface is about 140 nm or less.

9. The surface mechanical attrition treatment process of claim 3, wherein the room temperature processed metallic system has a surface Vickers microhardness of about 110 kg/mm²or more at room temperature.

10. The surface mechanical attrition treatment process of claim 3, wherein the cryogenic temperature processed metallic system has a surface Vickers microhardness of about 150 kg/mm²or more at cryogenic temperature.

11. The surface mechanical attrition treatment process of claim 4, wherein the room temperature refined grain size of the metallic system at its surface is about 635 nm or less.

12. The surface mechanical attrition treatment process of claim 4, wherein the cryogenic refined grain size of the metallic system at its surface is about 350 nm or less.

13. The surface mechanical attrition treatment process of claim 4, wherein the room temperature processed metallic system has a surface Vickers microhardness of about 230 kg/mm²or more at room temperature.

14. The surface mechanical attrition treatment process of claim 4, wherein the cryogenic temperature processed metallic system has a surface Vickers microhardness of about 260 kg/mm²or more at room temperature.

15. The surface mechanical attrition treatment process of claim 5, wherein the room temperature refined grain size of the metallic system at its surface is about 155 nm or less.

16. The surface mechanical attrition treatment process of claim 5, wherein the cryogenic refined grain size of the metallic system at its surface is about 60 nm or less.

17. The surface mechanical attrition treatment process of claim 5, wherein the room temperature processed metallic system has a surface Vickers microhardness of about 280 kg/mm²or more at room temperature.

18. The surface mechanical attrition treatment process of claim 5, wherein the cryogenic temperature processed metallic system has a surface Vickers microhardness of about 320 kg/mm²or more at room temperature.

19. The surface mechanical attrition treatment process of claim 1, wherein the impact processing time (t_i) is at least about 5 minutes.

20. The surface mechanical attrition treatment process of claim 1, wherein the impact processing time (t_i) is at least about 10 minutes.

21. The surface mechanical attrition treatment process of claim 1, wherein the impact processing time (t_i) is at least about 20 minutes.

22. The surface mechanical attrition treatment process of claim 1, wherein the impact processing time (t_i) is at least about 40 minutes.

23. The surface mechanical attrition treatment process of claim 1, wherein the impact processing time (t_i) is at least about 60 minutes.

24. A method of modifying the surface of a metal part, the method comprising the steps of:

providing at least one metal part t;

providing a plurality of metal fragments wherein said metal fragments have a size that is significantly smaller than the size of the at least one metal part;

reducing the temperature of the at least one metal part and the plurality of metal fragments; and

impacting the at least one metal part with the plurality of fragments at a reduced temperature (T_r), for an impact processing time (t_i).