Patent Application
20220048138
The present invention relates to additive manufacturing and, in particular, to a method for improving the strength and ductility of brittle intermetallic alloys through additive manufacturing.
Additive Manufacturing (AM) has seen tremendous growth in recent decades, largely motivated by opportunities to accelerate next-generation manufacturing processes and novel materials development. Principal and most well-discussed among the advantages of layerwise AM processing is the opportunity to implement exceptional design freedom, allowing the production of complex component geometries that are unachievable with conventional manufacturing techniques. See W. J. Sames et al., Int. Mater. Rev. 61, 315 (2016); W. E. Frazier, J. Mater. Eng. Perform. 23, 1917 (2014); and D. Herzog et al., Acta Mater. 117, 371 (2016). Navigation of this broad design space, often termed as “Design for Additive Manufacturing”, has been discussed in detail in the literature at various part, material, and product levels, and include aspects of both material performance and processing constraints. See M. K. Thompson et al., CIRP Ann.—Manuf. Technol. 65, 737 (2016). For parts, design opportunities include those that integrate material and geometric features toward enhancing functionality and production efficiency, such as conformal cooling, topology optimization for high strength-to-weight ratio, and cost-effective customized components. See M. K. Thompson et al., CIRP Ann.—Manuf. Technol. 65, 737 (2016); M. Garibaldi et al., J. Mech. Des. 141, 071401 (2019); and A. D. Dressler et al., Addit. Manuf. 28, 692 (2019). For material, the design space largely takes advantage of customized materials and micro/macrostructures that enable new properties and functionality. Some examples include customized powders and feedstock, tailored surfaces and porosity, and functionally graded materials. See I. E. Anderson et al., Curr. Opin. Solid State Mater. Sci. 22, 8 (2018); and D. C. Hofmann et al., J. Mater. Res. 29, 1899 (2014). Finally, at the product level, the design space encompasses the control of part consolidation, embedded parts, and the elimination of production assemblies. While these examples of additive design clearly exploit the advantageous processing characteristics of near-net-shape AM processing, the material properties of the constructed component remain, in many cases, inferior to conventionally processed wrought material.
An exemplary material is binary Fe—50Co alloy, a brittle intermetallic alloy of interest to electromagnetic applications. Fe—50Co, referred to hereafter as Fe—Co, exhibits exceptional magnetic properties that include high permeability, low coercivity, and high saturation induction, making these alloys ideal for use in a variety of electromagnetic devices. See T. Sourmail, Prog. Mater. Sci. 50, 816 (2005); and R. S. Sundar et al., J. Mater. Res. 20, 1515 (2005). However, when conventionally produced, this alloy is characterized by prohibitively low strength (200-400 MPa) and near-zero ductility (maximum 4% strain-at-failure), which challenges conventional processing and limits the application potential, for example in the aerospace industry. See T. Sourmail, Prog. Mater. Sci. 50, 816 (2005); and R. S. Sundar and S. C. Deevi, Soft magnetic FeCo alloys: alloy development, processing, and properties (2005). The poor mechanical properties of Fe—Co are proposed to be the result of a transition from disordered BCC to ordered B2 phases, which requires cooling rates in excess of 103° C./s to suppress, impractical for most conventional manufacturing methods. See D. W. Clegg and R. A. Buckley, Met. Sci. J. 7, 48 (1973); A. B. Kustas et al., Addit. Manuf. 21, 41 (2018); and M. Rajkovic et al., Met. Sci. 15, 21 (1981). The binary Fe—Co alloy has such unfavorable mechanical properties that, despite its functional performance, has never been commercialized in bulk form. See T. Sourmail, Prog. Mater. Sci. 50, 816 (2005). Rather, modified alloy compositions based on the Fe—Co binary system are routinely utilized, for example with small additions of ternary additives such as V and Nb. See K. Kawahara, J. Mater. Sci. 18, 1709 (1983). However, these commercial Fe—Co alloys (e.g., Fe—Co—2V Hiperco®, Permendur®) exhibit comparably harder magnetic properties, such as lower permeability and higher coercivity, leading to higher energy loss and larger/heavier electromagnetic devices. See T. Sourmail, Prog. Mater. Sci. 50, 816 (2005); and R. S. Sundar and S. C. Deevi, Soft magnetic FeCo alloys: alloy development, processing, and properties (2005). Substantial effort has been focused on developing techniques to process brittle magnetic alloys with optimal composition, however, they are largely limited to small volumes of material that remain brittle. See C. C. Lima et al., J. Alloys Compd. 586, S314 (2014); N. E. Fenineche et al., Mater. Lett. 58, 1797 (2004); Z. Turgut et al., J. Appl. Phys. 83, 6468 (1998); J. Xie et al., Intermetallics. 23, 20 (2012); E. J. Yun et al., IEEE Trans. Magn. 32, 4535 (1996); C. L. Platt et al., J. Appl. Phys. 88, 2058 (2000); G. Tian and X. Bi, J. Alloys Compd. 502, 1 (2010); Y. Shimada and H. Kojima, J. Appl. Phys. 47, 4156 (1976); R. Li et al., J. Magn. Magn. Mater. 281, 135 (2004); H. Haiji et al., J. Magn. Magn. Mater. 160 (1996); Y. Takada et al., J. Appl. Phys. 64, 5367 (1988); M. Abe et al., J. Mater. Eng. 11, 109 (1989); M. Abdellaoui et al., J. Alloys Compd. 198, 155 (1983); C. Kuhrt and L. Schultz, J. Appl. Phys. 73, 6588 (1993); A. B. Kustas et al., J. Mater. Process. Technol. 257 (2018); A. B. Kustas et al., Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 47, 3095 (2016); and B. Zhang et al., J. Magn. Magn. Mater. 324, 495 (2012). Emerging efforts in 3D printing of soft magnetic alloys with novel compositions are aimed toward producing components with unprecedented energy conversion efficiency. See G. Rolink et al., J. Mater. Res. 29, 2036 (2014); M. Garibaldi et al., Acta Mater. 110, 207 (2016); M. Garibaldi et al., Scr. Mater. 142, 121 (2018); M. Garibaldi et al., Mater. Charact. 1 (2018); C. V. Mikler et al., Jom. 69, 532 (2017); C. V. Mikler et al., Mater. Lett. 199, 88 (2017); C. V. Mikler et al., Mater. Lett. 192, 9 (2017); J. Geng et al., Jom. 68, 1972 (2016); A. Plotkowski et al., Addit. Manuf., 100781 (2019); A. B. Kustas et al., Addit. Manuf. 28, 772 (2019); and T. Riipinen et al., Rapid Prototyp. J. 25, 699 (2019).
Thus, while the design opportunities enabled by AM processing may promote improved performance margins relative to conventionally manufactured geometries, there remains a need for enhancing material performance for a given set of geometric and environmental constraints via control of the part construction.
The present invention is directed to a laser powder bed fusion (LPBF) AM processing method that enables drastically improved performance in a traditionally low strength, low ductility components and material systems. The method improves the general manufacturability of a component by overcoming inherent constraints associated with the process. Specifically, an arbitrary component shape can be produced via LPBF concomitantly with removable secondary artifacts, i.e., “heat sink” struts, to tailor the final component properties without modifying the final end-use geometry or part orientation. In particular, thermal profile of the component can be modified to keep the temperature of the intermetallic alloy below an alloy ordering temperature in order to keep the atomic mobility low and avoid excessive chemical ordering. For example, the method can be used with alloys that have a disordered BCC lattice that can transform to the B2 phase, whereby this transformation is suppressed. Intermetallic alloys that exhibit this transformation include, for example, Fe—Co, Fe—Al, and Fe—Si.
The mechanical properties of the LPBF process-modified specimens greatly exceed conventionally processed and LPBF-processed material without the secondary features, in some cases by up to an order-of-magnitude. Further, the method to improve mechanical properties can be used with other solidification-based additive processes when applied to traditionally brittle intermetallic alloys having disorder-order phase transformations.
As an example of the invention, a novel removable heat sink artifact design, uniquely enabled by metal additive manufacturing, was shown to drastically improve the performance of a traditionally low strength and low ductility Fe—Co alloy. The combination of yield strength and ductility of the AM-processed material was unprecedented, with a 300% increase in strength and an order-of-magnitude improvement in ductility relative to conventional wrought material. Thermal AM process simulations indicated that the design without heat sinks reached a much higher temperature during processing for a prolonged time, likely leading to a higher degree of atomic ordering. The processing temperatures seen in the simulations for designs with heat sinks were generally much lower, suggesting increased atomic disorder. Production and testing/use yield were also shown to drastically improve with the addition of the heat sink artifacts.
By utilizing a combination of “Design for Additive Manufacturing” principles with detailed fundamental materials understanding, revolutionary improvements in component performance can be realized. The designed processing methodology, unique to layerwise AM, is enables the development of higher performance alloys and net-shape devices. In terms of potential application impact, aerospace technologies, such as power generating units, internal generators for propulsion engines, and magnetic bearings, would benefit from the use of high strength and high ductility soft magnetic alloys.
The detailed description will refer to the following drawings, wherein like elements are referred to by like numbers.
As shown in
To demonstrate the invention, a conventional ASTM-based tensile specimen with a cylinder-based gauge section was selected as an exemplary geometry for manufacturing via LPBF, as shown in
The three tensile specimen designs were processed from binary Fe-50Co pre-alloyed powder (Sandvik) with a 15-45 μm distribution using a Renishaw AM400 (Wotton-under-Edge, UK) LPBF machine. Tensile specimens had dimensions of a 4 mm gauge diameter with a 22 mm gauge length and M12 printed threads on each end of the grips. The total length of each specimen was 54 mm. To isolate the role of the process design, processing conditions were held constant with a laser power of 200 W and an effective scan velocity of 0.833 m/s. An expanded set of processing conditions was explored for the V1 design by modifying the laser power over a range of 200 to 300 W, and effective scan velocities of 0.417 to 0.833 m/s to impose varying energy densities and investigate its effect on mechanical performance. Some of these additional processing conditions were not achievable for the specimen without heat sinks due to excessive heat accumulation and subsequent failure/drooping of the threaded ends during construction. All specimens were tested in the as-built condition after being cut off the build plate and without any secondary processing. Samples were produced using the “Meander” Renishaw scan strategy and a minimum of three specimens were tested for each design.
The effect of strut additions on thermal history was also evaluated using AM process simulations with the Sierra/Aria thermal finite element software suite. See S. T. D. Team, SAND2019-12296 (2019). This software package has been previously used to model the AM process, and a full description of the numerical formulation can be found in Johnson. See K. L. Johnson et al., Comput. Mech. 61, 559 (2018). These simulation differ slightly from Johnson due to the use of a planar heat source to activate a layer of material at a time rather than a Gaussian heat source matching the physical laser size. This change was made primarily due to the prohibitive computational expense of simulating each laser pass for a large LPBF part. Using a planar heat source sacrifices local thermal differences due to the scan path but is computationally efficient and well-suited to demonstrate differing thermal trends caused by geometry changes; e.g., the addition of struts.
A single build of each of the three different part geometries was modeled separately and centered on a 140×140×15 mm stainless steel baseplate representative of the LPBF process. The specimens were all meshed using tetrahedral elements and a desired element size of 0.3 mm, while the requested element size for the baseplate was 4 mm. The transition from fine to coarse elements occurred in the baseplate. A planar heat source with 0.25 mm thickness representing one agglomerated layer (equal to five physical layers) was translated from the baseplate to the top of the part. At each 0.25 mm layer location, the source was held for 0.5 s and then deactivated for 110 s before moving to the next layer. The 110 s delay was chosen in an attempt to replicate the correct cooling time for the experimental parts. The Renishaw machine used was operated in reduced build volume mode, which requires approximately 22 s for recoating time per physical layer. Because the physical layer thickness was 0.05 mm, an agglomerated layer thickness of 0.25 mm would require 5× the cooling time or 110 s. The thermophysical properties used in the simulations are shown in Table 1 below. In addition to those in the table, the latent heat of fusion value of 246.6 kJ/kg was applied across a 25-degree temperature range beginning at the solidus temperature of 1750K. See T. G. Woodcock et al., Calphad Comput. Coupling Phase Diagrams Thermochem. 31, 256 (2007); and J. E. Rodriguez et al., Acta Mater. 122, 431 (2017). Convection was only considered on the exterior surfaces of the parts and baseplate. To mimic the effect of surrounding powder, a convection coefficient of 8 W/m2·K was considered on the part exterior, while a value of 25 W/m2·K was used on the bottom of the baseplate to mimic partial conduction to the LPBF mounting brackets. See A. J. Dunbar, Analysis of the laser powder bed fusion additive manufacturing process through experimental measurement and finite element modeling (2016).
1Carpenter Hiperco 50 Fe-Co-V Soft Magnetic Alloy, 0.15 mm Strip Heat Treated 843 C.
2P. Choufani, Measuring the Thermophysical Properties of Fe-Co alloys using Containerless Electrostatic Levitation Techniques (2012).
3G. Kozlowski et al., Contin. Mech. Thermodyn. 32, 247 (2020).
4A. S. Normanton et al., Met. Sci. 9, 510 (1975).
The resultant engineering tensile stress vs. strain curves for the three different design approaches are shown in
The slight deterioration of the mechanical properties for the V2 specimen, as compared to V1, could be from increased surface defects along the gauge length that resulted during the removal of the larger number of struts prior to tensile testing. This may reduce ductility when considering that FeCo-based alloys generally have a high notch sensitivity. See N. Stoloff, Ordered Alloys for High Temperature Applications, MRS Proc. 39 (1985); and M. R. Pinnel et al., Acta Metall. 24, 1095 (1976). Additionally, microstructural analysis revealed that the V1 specimens may, on average, have had fewer AM-induced defects compared to the V2 and V0 specimens. This is evident in
One possible reason for the differences in the defect density between the three specimens is that the defect variation resulted from the general stochastic nature of AM processing. Other possibilities may be linked to the different temperature profiles experienced by each design. For the V2 specimens, it is conceivable that the increased number of heat sinks could have generated a temperature profile that was below the alloy ductile-to-brittle transition temperature (˜450° C.), promoting defects during processing due to poor alloy workability. See T. L. Johnston et al., Philos. Mag. 12, 305 (1965); and N. S. Stoloff and R. G. Davies, Prog. Mater. Sci. 13, 1 (1968). For the V0 specimen, the temperatures during processing may have been high enough to transition to an austenitic microstructure. Repetitively cycling between austenite and ferrite phases could have generated additional residual stress, leading to an increase in defects. Consequentially, the V0 specimens may have been processed with a temperature profile between these two extremes, leading to near defect-free microstructures. Finally, slight differences in grain structure may have also affected the tensile performance (see
Quantitative analysis of the tensile response for the three AM-produced tensile specimens, with a comparison to literature data for conventionally processed Fe—Co, is shown in
The AM-processed specimens also had, in some cases, significantly higher ductility compared to conventional material, as shown in
While the impact of the designed processing methodology is apparent in enabling unusual performance for exemplar tensile specimens, the role of processing conditions was also important to consider.
To contextualize the impact of the processing method of the invention, the product performance, production yield, and manufacturing implications are considered. For the former, consider the yield strength vs. total strain plot in
The thermal simulations further confirm the reduction in temperature rise for the specimens with heat sinks. Temperature histories were recorded at a point in the center of the gauge section for each of the specimen geometries, as shown in
Soon after the heat source reached the point of interest, the thermal histories were very similar, with V2 showing slightly lower temperatures overall due to the numerous struts attaching throughout the length of the specimen. The most dramatic differences occurred once the heat source reached the grip sections. Due to the large increase in the cross-sectional area for the grip, as compared to the gauge section, excessive heat accumulated in the part. This led to a significant temperature rise for the grip far above what was experienced by the gauge section. At approximately 18,000 s, the V1 struts, which were previously disconnected from the tensile specimen, joined the grip section and led to a significant decrease in temperature. Conversely, for the geometry with no heat sinks, the temperature profile continued to increase. Perhaps as expected, V2 showed only a minor increase in temperature during this stage of the build due to the repeating strut attachments to the tensile specimen. At approximately 22,000 s, the temperatures dropped slightly due to an internal hex key-shaped channel in the top of the grip that reduced the local cross-sectional area. After completion of the build, the heat source was removed and the parts cooled down to ambient temperature.
Though these temperature profiles are qualitative, they give important insights into a likely mechanism for the different mechanical responses between designs. The thermal predictions suggest that the V0 specimen experienced temperatures above the alloy ordering temperature, 730° C. (1003K), and possibly into the austenitic phase-field (983° C., 1256K) for nearly an hour (19,000 to 22,000 s). Under these types of thermal conditions, the material would likely cycle between multiple phase regions. Notably, the temperature profile for the V0 specimens near the end of the build (from ˜22,000 s to end of the simulation) was in a high atomic mobility region for the B2 phase transformation. This suggests that the V0 design had a significant fraction of the B2 phase, which would reduce strength and ductility. See D. W. Clegg and R. A. Buckley, Met. Sci. J. 7, 48 (1973). Conversely, the V1 and V2 specimens were likely heated several hundreds of degrees below the disorder-order phase boundary in a region of lower atomic mobility, limiting the extent of the B2 phase transformation and enabling higher strength and ductility.
At the material level, mechanisms leading to the drastically improved mechanical properties for the Fe—Co alloy are speculative. However, the rapid cooling and minimal layer-by-layer temperature rise facilitated by AM has been shown to promote the formation of an atomically disordered BCC structure in Fe—Co. See A. B. Kustas et al., Addit. Manuf. 21, 41 (2018); and A. B. Kustas et al., Addit. Manuf. 28, 772 (2019). These thermal conditions were likely enhanced by the local strut artifact design, encouraging suppression of the B2 ordered phase, and promoting higher strength and higher ductility. See T. Sourmail, Prog. Mater. Sci. 50, 816 (2005); and R. S. Sundar and S. C. Deevi, Soft magnetic FeCo alloys: alloy development, processing, and properties (2005). Thus, it is possible that local thermal management via AM processing was the sole mechanism leading to enhanced properties for the Fe—Co alloy through stabilization of the disordered BCC structure. However, the performance of AM-processed material was significantly improved relative to that typically observed in fully disordered material when conventionally processed. Therefore, additional atomic-scale mechanisms, inherited from the AM processing, must be present in order to facilitate the remarkable properties. This is especially true considering that AM-produced Fe—Co can exhibit relatively poor mechanical properties in some instances, as described herein for the higher energy density processing conditions and in other recent evaluations. See T. Riipinen et al., Rapid Prototyp. J. 25, 699 (2019); and T. F. Babuska et al., Acta Mater. 180, 149 (2019). A tentative hypothesis, based on high-resolution material characterization and atomic-scale computational analysis, suggests that rapid solidification of LPBF-type AM processes can be utilized to produce microstructural features that act to simultaneously impede and accommodate dislocations during plastic deformation. See T. F. Babuska et al., Acta Mater.180, 149 (2019). Microstructural features, such as nanoscale voids, dislocation cells, sub-micron cellular/dendritic structures, oxides, etc., have been noted to form in LPBF-processed material, including Fe—Co. These features could enable additional plastic deformation accommodation mechanisms and lead to higher strength and extended ductility, as was similarly proposed by Wang for AM-processed 304L stainless steel. See Y. M. Wang et al., Nat. Mater. 17, 63 (2018). The formation of these features would likely be enhanced by the application of local heat sink artifacts integrated for augmented heat transfer and cooling rates. In addition to the multiscale microstructural features typical of AM-processed material, rapid cooling, similar to the cooling rates found in laser-based AM processing, has been shown to promote martensitic phases in Fe—Co—V alloys, which may influence mechanical behavior. See S. Mahajan et al., Met. Trans. 5, 1263 (1974); and J. A. Ashby et al., Met. Sci. 11, 91 (1977). However, no such phases have been confirmed in AM-processed binary Fe—Co to date. Finally, the AM-processed Fe—Co alloy specimens all possessed a relatively fine (5-20 μm diameter), predominately equiaxed grain structure in the as-built condition (see
A practical processing benefit of the process design methodology was also identified when comparing the production yield/efficiencies for the various specimen designs. In the case of the design without heat sinks, 50% of the specimens were constructed with sufficient quality to enable mechanical testing/use. Specimens that could not be tested generally failed during the layerwise processing from excessive heat accumulation and produced low quality threaded ends that were unusable with the tensile grips. This result is consistent with the simulation results, which predicted near melt temperatures in the grips for the V0 specimens. Product and testing yield were found to significantly increase for specimens with the removable heat sink artifacts for both V1 and V2 specimens, which yielded 75% and 90% test/use rates, respectively, see Table 3 for details. These findings demonstrate the value of utilizing AM to design for improved manufacturability to increase the rate of component use in line with traditional “Design for Additive Manufacturing” guidelines.
As shown in
The present invention has been described as a method of improving the strength and ductility of brittle intermetallic alloys through additive manufacturing. It will be understood that the above description is merely illustrative of the applications of the principles of the present invention, the scope of which is to be determined by the claims viewed in light of the specification. Other variants and modifications of the invention will be apparent to those of skill in the art.
This invention was made with Government support under Contract No. DE-NA0003525 awarded by the United States Department of Energy/National Nuclear Security Administration. The Government has certain rights in the invention.
