Patent Grant
8882941
The present invention relates the formation of spinodal microconstituent structures in a metallic glass matrix which exhibit combinations of relatively high tensile strength and relatively high elongation
Metallic nanocrystalline materials and metallic glasses exhibit relatively high hardness and strength characteristics for metal-based materials and because of this, they are considered to be potential candidates for structural applications. However, their limited fracture toughness and ductility associated with the rapid propagation of shear bands and/or cracks may be a concern for the commercial utilization of their superior strength. Typically, these materials may exhibit adequate ductility by testing in compression while tensile ductility of the same materials may be close to zero. At the same time, tensile ductility along with fracture toughness is understood to be a relatively important characteristic for structural applications where intrinsic ductility is needed to avoid catastrophic failure.
Nanocrystalline materials may be understood to be or include polycrystalline structures with a mean grain size below 100 nm. They have been the subject of widespread research since mid-1980s when it was asserted that metals and alloys, if made nanocrystalline, may exhibit a number of appealing mechanical characteristics of potential significance for structural applications. But despite relatively attractive properties (high hardness, yield stress and fracture strength), it is understood that they may show a disappointingly low tensile elongation and may tend to fail in a relatively brittle manner. In fact, empirical correlations between the work hardening exponent and the grain size for cold rolled and conventionally recystallized mild steels indicate a decrease in ductility for decreasing grain size. As the grain size is progressively decreased, the formation of dislocation pile-ups may become more difficult, limiting the capacity for strain hardening. That may lead to mechanical instability of materials under loading.
Attempts to improve the ductility of nanocrystalline materials while maintaining their relatively high strength by adjusting the microstructure have also been made. It has been proposed that an increased content of high-angle grain boundaries in nanocrystalline materials could be beneficial to an increase in ductility. In a search to improve ductility of nanocrystalline materials, extremely ductile base metals have been used. For example, nanocrystalline Cu with a bimodal grain size distribution (100 nm and 1.7 μm) has been fabricated based on the thermomechanical treatment of severe plastic deformation, which may exhibit a 65% total elongation to failure and may retain a relative high strength. Recently, a nanocrystalline Cu with nanometer sized twins embedded in submicrometer grained matrix by pulsed electrodepositon has been produced. The ductility and relatively high strength may be attributed to the interaction of glide dislocations with twin boundaries. In a recent approach, nanocrystalline second-phase particles of 4-10 nm were incorporated into the nanocrystalline Al matrix (about 100 nm). The nanocrystalline particles interacted with the slipping dislocation and enhanced the strain hardening rate which leads to the evident improvement of ductility. Using these approaches, enhanced tensile ductility has been achieved in a number of nanocrystalline materials such as 15% in pure Cu with mean grain size of 23 nm or 30% in pure Zn with mean grain size of 59 nm. It should be noted that fracture strength of these nanocrystalline materials does not exceed 1 GPa. For nanocrystalline materials with higher fracture strength (1 GPa) the achievement of adequate ductility (>1%) may still be a challenge.
Amorphous metallic alloys (metallic glasses) represent a relatively young class of materials, having been first reported around 1960 when classic rapid-quenched experiments were performed on Au—Si alloys. Since that time, there has been progress in exploring alloys compositions for glass formers, seeking elemental combinations with ever-lower critical cooling rates for the retention of an amorphous structure. Due to the absence of long-range order, metallic glasses may exhibit relatively unique properties, such as relatively high strength, high hardness, large elastic limit, good soft magnetic properties and high corrosion resistance. However, owing to strain softening and/or thermal softening, plastic deformation of metallic glasses may be highly localized into shear bands, which may result in a limited plastic strain (less than 2%) and may lead to catastrophic failure at room temperature.
Different approaches have been applied to enhance ductility of metallic glasses such as introducing free volume in amorphous structure or glassy phase separation which has enabled up to 25% in compression. However, tensile ductility for these materials has not been reported. Another approach is the development of metallic glass matrix composites. Crystalline precipitates may be introduced into a glass matrix by partial crystallization. Crystallization occurs by nucleation and growth mechanism and depending on glass composition and crystallization kinetics, nanometer-sized or micrometer-sized crystallities might be introduced.
This approach may also allow an increase in compressive ductility in Ti-based, Zr-based, Mg-based glasses and Cu—Hf—Ti—Nb system. Tensile ductility of these materials was demonstrated up to 13% tensile elongation in Ti—Zr-based metallic glasses with large dendrites (20-50 μm in size) embedded in the glassy matrix. The heterogeneous structure of these composites may act as an initiation site for the formation of shear bands and/or a barrier to the rapid propagation of shear bands, which leads to enhancement of global plasticity, but sometimes decreases the strength.
Another way to reduce grain size is through spinodal decomposition which may occur when a mixture of two or more materials separate into distinct regions with different material concentrations. This method differs from nucleation in that phase separation due to spinodal decomposition may occur throughout the material, and not just at nucleation sites. Spinodal decomposition was previously observed in AlNiCo magnets, 17-4PH stainless steel, Fe-25Cr-12Co-1Si alloy, and Fe-based austenitic alloy. Recent studies mentioned a Co enrichment in the amorphous residual matrix and Fe enrichment in the α′-FeCo crystalline phase. In addition, experimental evidence of grain refinement caused by the formation of clusters, which, in turn, resulted from the addition of >1% Cu has been presented. It has also been shown that Cu additions of more than 1.0% promoted the formation of clusters responsible for the grain refinement of the crystalline α′-FeCo phase, i.e., grain size around 10 nm. However, no property evaluation of the final structure was performed in these studies. With respect to AlNiCo magnets, while it is known from a number of sources that relatively high tensile strengths can be obtained from 28 to 380 MPa, the material response may be somewhat brittle and tensile elongation data is generally not listed.
An aspect of the present disclosure relates to an alloy composition, which may include 52 atomic percent to 68 atomic percent iron, 13 to 21 atomic percent nickel, 2 to 12 atomic percent cobalt, 10 to 19 atomic percent boron, optionally 1 to 5 atomic percent carbon, and optionally 0.3 to 16 atomic percent silicon. The alloy may include 5 to 95% by volume of one or more spinodal microconstituents, wherein the microconstituents exhibit a length scale less than 50 nm in a glass matrix.
Another aspect of the present disclosure relates to a method of forming spinodal microconstituents in an alloy. The method may include melting alloy constituents including 52 atomic percent to 60 atomic percent iron, 15.5 to 21 atomic percent nickel, 6.3 to 11.6 atomic percent cobalt, 10.3 to 13.2 atomic percent boron, 3.7 to 4.8 atomic percent carbon, and 0.3 to 0.5 atomic percent silicon to form an alloy, and cooling the alloy to form one or more spinodal microconstituents in a glass matrix. The spinodal microconstituents may be present in the range of 5% to 95% by volume and exhibit a length scale less than 50 nm in a glass matrix.
The above-mentioned and other features of this disclosure, and the manner of attaining them, may become more apparent and better understood by reference to the following description of embodiments described herein taken in conjunction with the accompanying drawings, wherein:
The present disclosure relates to a glass forming alloy which may transform to yield at least a portion of its structure as a spinodal microconstituent, which may consist of one or more crystalline phases at a length scale less than 50 nm in a glass matrix. Stated another way, any given dimension of the crystalline phases may be in the range of 1 nm to less than 50 nm including all values and increments therein, such as 1 nm, 2 nm, 3 nm 4 nm, 5 nm, 6 nm, 7 nm, 8 nm, 9 nm, 10 nm, 11 nm, 12 nm, 13 nm, 14 nm, 15 nm, 16 nm, 17 nm, 18 nm, 19 nm, 20 nm, 21 nm 22 nm 23 nm, 24 nm, 25 nm, 26 nm, 27 nm, 28 nm, 29 nm, 30 nm, 40 nm, 41 nm, 42 nm, 43 nm, 44 nm, 45 nm, 46 nm, 47 nm, 48 nm, 49 nm. In addition, the alloy may include one or more of spinodal microconstituents present in the range of ˜5 to ˜95% by volume, including 5%, 6%, 7%, 8%, 9%, 10%, 11%, 12%, 13%, 14%, 15%, 16%, 17%, 18%, 19%, 20%, 21%, 22%, 23%, 24%, 25%, 26%, 27%, 28%, 29%, 30%, 31%, 32%, 33%, 34%, 35%, 36%, 37%, 38%, 39%, 40%, 41%, 42%, 43%, 44%, 45%, 46%, 47%, 48%, 49%, 50%, 51%, 52%, 53%, 54%, 55%, 56%, 57%, 58%, 59%, 60%, 61%, 62%, 63%, 64%, 65%, 66%, 67%, 68%, 69%, 70%, 71%, 72%, 73%, 74%, 75%, 76%, 77%, 78%, 79%, 80%, 81%, 82%, 83%, 84%, 85%, 86%, 87%, 88%, 89%, 90%, 91%, 92%, 93%, 94%, 95%. Spinodal microconstituents may be understood as microconstituents formed by a transformation mechanism which is not nucleation controlled. More basically, spinodal decomposition may be understood as a mechanism by which a solution of two or more components (e.g. metal compositions) of the alloy can separate into distinct regions (or phases) with distinctly different chemical compositions and physical properties. This mechanism differs from classical nucleation in that phase separation occurs uniformly throughout the material and not just at discrete nucleation sites. One or more semicrystalline clusters or crystalline phases may therefore form through a successive diffusion of atoms on a local level until the chemistry fluctuations lead to at least one distinct crystalline phase. Semi-crystalline clusters may be understood herein as exhibiting a largest linear dimension of 2 nm or less, whereas crystalline clusters may exhibit a largest linear dimension of greater than 2 nm. Note that during the early stages of the spinodal decomposition, the clusters which are formed are small and while their chemistry differs from the glass matrix, they are not yet fully crystalline and have not yet achieved well ordered crystalline periodicity. Additional crystalline phases may exhibit the same crystal structure or distinct structures.
Glass forming alloys that may provide spinodal microconstituent formation may include the following constituents: 52 to 68 atomic percent (at %) iron, 13 to 21 at % nickel, 2 to 12 at % cobalt, 10 to 19 at % boron, 1 to 5 at % carbon if present, 0.3 to 16 at % silicon if present, including all values and increments of 0.1 atomic percent within the above ranges. For example, the glass forming alloys may include 52 atomic percent to 60 atomic percent iron 15.5 to 21 atomic percent nickel, 6.3 to 11.6 atomic percent cobalt, 10.3 to 13.2 atomic percent boron, 3.7 to 4.8 atomic percent carbon; and 0.3 to 0.5 atomic percent silicon. In another example, the glass forming alloys may include 58.4 atomic percent to 67.6 atomic percent iron, 16.0 to 16.6 atomic percent nickel, 2.9 to 3.1 atomic percent cobalt, 12.0 to 18.5 atomic percent boron, optionally 1.5 to 4.6 atomic percent carbon, and optionally 0.4 to 3.5 atomic percent silicon. In yet a further example, the glass forming alloys may include 53.6 atomic percent to 60.9 atomic percent iron, 13.6 to 15.5 atomic percent nickel, 2.4 to 2.9 atomic percent cobalt, 12 to 14.1 atomic percent boron, 1 to 4 atomic percent carbon, and 3.9 to 15.4 atomic percent silicon. It may be appreciated that the alloys may not only include, but may also consist essentially of or consist of the above described constituents. Furthermore, even where the alloys consist of the above, it may be appreciated that some degree of impurities may be present in the alloy compositions, such as in the range of 0.01 to 1.0 atomic percent of impurities, including all values and increments therein at 0.01 atomic percent increments.
Accordingly, it may therefore be appreciated that iron may be present at 52.0, 52.1, 52.2, 52.3, 52.4, 52.5, 52.6, 52.7, 52.8, 52.9, 53.0, 53.1, 53.2, 53.3, 53.4, 53.5, 53.6, 53.7, 53.8, 53.9, 54.0, 54.1, 54.2, 54.3, 54.4, 54.5, 54.6, 54.7, 54.8, 54.9, 55.0, 55.1, 55.2, 55.3, 55.4, 55.5, 55.6, 55.7, 55.8, 55.9, 56.0, 56.1, 56.2, 56.3, 56.4, 56.5, 56.6, 56.7, 56.8, 56.9, 57.0, 57.1, 57.2, 57.3, 57.4, 57.5, 57.6, 57.7, 57.8, 57.9, 58.0, 58.1, 58.2, 58.3, 58.4, 58.5, 58.6, 58.7, 58.8, 58.9, 59.0, 59.1, 59.2, 59.3, 59.4, 59.5, 59.6, 59.7, 59.8, 59.9, 60.0, 60.1, 60.2, 60.3, 60.4, 60.5, 60.6, 60.7, 60.8, 60.9, 61.0, 61.1, 61.2, 61.3, 61.4, 61.5, 61.6, 61.7, 61.8, 61.9, 62.0, 62.1, 62.2, 62.3, 62.4, 62.5, 62.6, 62.7, 62.8, 62.9, 63.0, 63.1, 63.2, 63.3, 63.4, 63.5, 63.6, 63.7, 63.8, 63.9, 64.0, 64.1, 64.2, 64.3, 64.4, 64.5, 64.6, 64.7, 64.8, 64.9, 65.0, 65.1, 65.2, 65.3, 65.4, 65.5, 65.6, 65.7, 65.8, 65.9, 66.0, 66.1, 66.2, 66.3, 66.4, 66.5, 66.6, 66.7, 66.8, 66.9, 67.0, 67.1, 67.2, 67.3, 67.4, 67.5, 67.6, 67.7, 67.8, 67.9, 68.0 atomic percent. It may also be appreciated that nickel may be present at 13.1, 13.2, 13.3, 13.4, 13.5, 13.6, 13.7, 13.8, 13.9, 14.0, 14.1, 14.2, 14.3, 14.4, 14.5, 14.6, 14.7, 14.8, 14.9, 15.0, 15.1, 15.2, 15.3, 15.4, 15.5, 15.6, 15.7, 15.8, 15.9, 16.0, 16.1, 16.2, 16.3, 16.4, 16.5, 16.6, 16.7, 16.8, 16.9, 17.0, 17.1, 17.2, 17.3, 17.4, 17.5, 17.6, 17.7, 17.8, 17.9, 18.0, 18.1, 18.2, 18.3, 18.4, 18.5, 18.6, 18.7, 18.8, 18.9, 19.0, 19.1, 19.2, 19.3, 19.4, 19.5, 19.6, 19.7, 19.8, 19.9, 20.0, 20.1, 20.2, 20.3, 20.4, 20.5, 20.6, 20.7, 20.8, 20.9, 21.0 atomic percent. Cobalt may be present at 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6, 5.7, 5.8, 5.9, 6.0, 6.1, 6.2, 6.3, 6.4, 6.5, 6.6, 6.7, 6.8, 6.9, 7.0, 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9, 8.0, 8.1, 8.2, 8.3, 8.4, 8.5, 8.6, 8.7, 8.8, 8.9, 9.0, 9.1, 9.2, 9.3, 9.4, 9.5, 9.6, 9.7, 9.8, 9.9, 10.0, 10.1, 10.2, 10.3, 10.4, 10.5, 10.6, 10.7, 10.8, 10.9, 11.0, 11.1, 11.2, 11.3, 11.4, 11.5, 11.6, 11.7, 11.8, 11.9, 12.0 atomic percent. Boron may be present at 10.0, 10.1, 10.2, 10.3, 10.4, 10.5, 10.6, 10.7, 10.8, 10.9, 11.0, 11.1, 11.2, 11.3, 11.4, 11.5, 11.6, 11.7, 11.8, 11.9, 12.0, 12.1, 12.2, 12.3, 12.4, 12.5, 12.6, 12.7, 12.8, 12.9, 13.0, 13.1, 13.2, 13.3, 13.4, 13.5, 13.6, 13.7, 13.8, 13.9, 14.0, 14.1, 14.2, 14.3, 14.4, 14.5, 14.6, 14.7, 14.8, 14.9, 15.0, 15.1, 15.2, 15.3, 15.4, 15.5, 15.6, 15.7, 15.8, 15.9, 16.0, 16.1, 16.2, 16.3, 16.4, 16.5, 16.6, 16.7, 16.8, 16.9, 17.0, 17.1, 17.2, 17.3, 17.4, 17.5, 17.6, 17.7, 17.8, 17.9, 18.0, 18.1, 18.2, 18.3, 18.4, 18.5, 18.6, 18.7, 18.8, 18.9, 19.0 atomic percent. Carbon may be present at 0.0, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0 atomic percent. Silicon may be present at 0.0, 0.3, 0.4, 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6, 5.7, 5.8, 5.9, 6.0, 6.1, 6.2, 6.3, 6.4, 6.5, 6.6, 6.7, 6.8, 6.9, 7.0, 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9, 8.0, 8.1, 8.2, 8.3, 8.4, 8.5, 8.6, 8.7, 8.8, 8.9, 9.0, 9.1, 9.2, 9.3, 9.4, 9.5, 9.6, 9.7, 9.8, 9.9, 10.0, 10.1, 10.2, 10.3, 10.4, 10.5, 10.6, 10.7, 10.8, 10.9, 11.0, 11.1, 11.2, 11.3, 11.4, 11.5, 11.6, 11.7, 11.8, 11.9, 12.0, 12.1, 12.2, 12.3, 12.4, 12.5, 12.6, 12.7, 12.8, 12.9, 13.0, 13.1, 13.2, 13.3, 13.4, 13.5, 13.6, 13.7, 13.8, 13.9, 14.0, 14.1, 14.2, 14.3, 14.4, 14.5, 14.6, 14.7, 14.8, 14.9, 15.0, 15.1, 15.2, 15.3, 15.4, 15.5, 15.6, 15.7, 15.8, 15.9, 16.0 atomic percent.
The alloys may also exhibit a critical cooling rate for metallic glass formation of about <100,000 K/s. Critical cooling rate may be understood as a rate of continuous cooling which may suppress and/or reduce transformations, which may be undesirable, such as crystallization. Accordingly, the alloys may be formed by melting and cooling the alloys at or below the critical cooling rate avoiding glass devitrification and forming a supersaturated matrix. The supersaturated matrix may then undergo spinodal decomposition forming spinodal microconstituents. Methods of forming the alloys (including melting and/or cooling the alloys) include those methods that may allow for the alloys to cool at a rate that is equal to or greater than the critical cooling rate, such as melt spinning. In addition, the alloy may be processed to yield a thin product from 1 μm to 2000 μm in thickness in the form of a powder particle, thin film, flake, ribbon, wire, or sheet. An example of an alloy forming technique may include melt spinning, jet casting, Taylor-Ulitovsky, melt-overflow, planar flow casting, and twin roll casting.
The alloy may exhibit a density in the range of 7 to 8 grams per cubic centimeter, including all values and increments therein, as measured by the Archimedes method, such as 7.1, 7.2, 7.3, 7.4, 7.5, 7.6, 7.7, 7.8, 7.9 8.0 grams per cubic centimeter. The alloys may also exhibit one or more onset crystallization temperature in the range of 400° C. to 585° C., including all values and increments therein in 1° C. increments, measured by DTA at 10° C./min. The alloy may exhibit one or more a peak crystallization temperatures in the range of about 400 to 595° C., including all values and increments therein in 1° C. increments, measured by DTA at 10° C./min. In addition, the alloys may exhibit one or more onset melting temperatures in the range of 1050° C. to 1100° C., including all values and increments therein in 1° C. increments, measured by DTA at 10° C./min and one or more peak melting temperature in the range of 1050° C. to 1125° C., including all values and increments therein in 1° C. increments. It can be appreciated that the onset temperatures occur before the respective peak temperatures and that multiple onset and peak crystallization and/or melting temperatures may be present.
The resulting microstructure of the alloys after being produced may therefore all include as a portion thereof a spinodal microconstituent which includes one or more crystalline phases uniformly dispersed at a length scale less than 50 nm. Reference to uniformly dispersed may be understood as noted above in that the spinodal microconstituent is formed via a phase separation that occurs within the sample material and not at discrete nucleation sites.
Such spinodal microstructure may also include all amorphous regions, isolated crystalline precipitates in a glass matrix, multiphase crystalline clusters growing into the glass matrix, completely crystalline areas with nanocrystalline crystallite from 10 to 100 nm, a three phase nanoscale microconstituent with about two relatively fine, i.e., less than 15 nm, including all values and increments in the range of 1 nm to 15 nm, crystalline phases intermixed in a glass matrix, as well as combinations thereof. In one example the resulting structure of the alloy may consist primarily of metallic glass. Reference to metallic glass may be understood as microstructures that may exhibit associations of structural units in the solid phase that may be randomly packed together. The level of refinement, or the size, of the structural units may be in the angstrom scale range (i.e. 5 Å to 100 Å).
In another example, the resulting structure of the alloys may consist of metallic glass and crystalline phases less than 500 nm in size, including all values and increments in the range of 10 nm to 500 nm in size. Furthermore, as noted above, the alloys may transform to yield at least a portion of its structure as a spinodal microconstituent which may consist of one or more crystalline phases at a length scale less than 50 nm in a glass matrix. In other words, the largest linear dimension of the semi-crystalline or crystalline phases may be in the range of 1 nm to 50 nm, including all values and increments therein.
The alloys may exhibit varying degrees of brittleness, and as measured by a bend test, i.e., bending of ribbons 180°, wherein the alloy samples could be bent on either side, on one side or could not bend without breaking. The alloy structure may exhibit a tensile elongation greater than 0.65%, including all values and increments in the range of 0.65% to 7.5% at 0.01 increments, such as 1 to 7.06%. In addition, the alloy may exhibit a yield strength greater than 0.1 GPa, including all values and increments in the range of 0.1 GPa to 2.2 GPa. The alloy may also exhibit an ultimate tensile strength of 0.1 GPa to 3.5 GPa, including all values and increments therein, a Young's Modulus of 55 GPa to 130 GPa, including all values and increments therein. The alloys herein are thus capable of providing one or more of the above referenced mechanical properties in combination.
The following examples are presented for purposes of illustration only and are not meant to limit the scope of the application. In addition, the examples may provide support for ranges within the specific points disclosed.
Sample Preparation
Using high purity elements (i.e., being 99% purity or greater), 15 g alloy feedstocks of the targeted alloys were weighed out according to the atomic ratio's provided in Table 1. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and remelted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a melt-spinning chamber in a quartz crucible with a hole diameter of ˜0.81 mm. The ingots were melted in a ⅓ atm helium atmosphere using RF induction and then ejected onto a 245 mm diameter copper wheel which was traveling at tangential velocities which typically were either 16 or 10.5 m/s. The resulting ribbons that were produced had widths which were typically ˜1.25 mm and thickness from 0.04 to 0.08 mm as shown in Table 2.
Density
The density of the alloys in ingot form was measured using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water. The density of the arc-melted 15 gram ingots for each alloy is tabulated in Table 3 and was found to vary from 7.73 g/cm3 to 7.85 g/cm3. Experimental results have revealed that the accuracy of this technique is +−0.01 g/cm3.
As-Solidified Structure
Thermal analysis was done on the as-solidified ribbon structure on a Perkin Elmer DTA-7 system with the DSC-7 option. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) was performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultrahigh purity argon. In Table 4, the DSC data related to the glass to crystalline transformation is shown for the alloys that have been melt-spun at two different wheel tangential velocities at 16 m/s and 10.5 m/s. Note that the cooling rate increases at increasing wheel tangential velocities. In
In Table 5, elevated temperature DTA results are shown indicating the melting behavior for the alloys shown in Table 1. As can be seen from the tabulated results in Table 4 and the melting peaks in
X-Ray Diffraction Analysis
The as-spun ribbons were cut into short segments and four to six pieces of ribbon were placed on an off-cut SiO2 single crystal (zero-background holder). The ribbons were situated such that either the shiny side (free side) or the dull side (wheel side) were positioned facing up on the holder. A small amount of silicon powder was placed on the holder as well, and then pressed down so that the height of the silicon matched the height of the ribbon, which allows for matching any peak position errors in subsequent detailed phase analysis. X-ray diffraction scans were taken from 20 to 100 degrees two theta with a step size of 0.02 degrees and at a scanning rate of 2 degrees/minute. The X-ray tube settings were measured with a copper target at 40 kV and 44 mA. In
TEM Analysis
Specimens for transmission electron microscopy (TEM) were produced from melt-spun ribbon by a combination of mechanical thinning and ion milling. The ribbons were mechanically thinned from their original thickness to approximately 10 microns using fine-grit sandpaper followed by polishing using 5 micron and 0.3 micron alumina powder on felt pads with water used as a lubricant in both cases. Ribbon sections of 3 mm were then cut using a razor blade and the resulting sections were mounted on copper support rings with two-part epoxy since the support rings provide structural integrity for handling. The specimens were then ion milled using a Gatan Precision Ion Polishing System (PIPS) operating at 4.5 kV. Incident angles were decreased from 9 degrees to 8 degrees and finally 7 degrees every ten minutes during the ion milling process. The resulting thin areas were examined using a JEOL 2010 TEM operating at 200 kV. For each alloy listed in Table 1, TEM micrographs were taken near the center of the ribbon thickness for samples melt-spun at both 16 m/s and 10 m/s. In
Mechanical Property Testing
Mechanical property testing was done primarily through using qualitative 180° bend testing and tensile testing. The following sections will detail the technical approach and measured data.
180 Degree Bend Testing
The ability of the ribbons to bend completely flat indicates a special condition whereby relatively high strain can be obtained but not measured by traditional bend testing. When the ribbons are folded completely around themselves, they experience relatively high strain which can be as high as 119.8% as derived from complex mechanics. In practice, the strain may be in the range of ˜57% to ˜97% strain in the tension side of the ribbon. During 180° bending (i.e. flat), four types of behavior were observed; Type 1 Behavior—not bendable without breaking, Type 2 Behavior—bendable on one side with wheel side out, Type 3 Behavior—bendable on one side with free side out, and Type 4 Behavior—bendable on both sides. In Table 6, a summary of the 180° bending results including the specific behavior type are shown for the studied alloys processed at both 16 and 10.5 m/s. In
Tensile Test Results
The mechanical properties of metallic ribbons were obtained at room temperature using microscale tensile testing. The testing was carried out in a commercial tensile stage made by Fullam which was monitored and controlled by a MTEST Windows software program. The deformation was applied by a stepping motor through the gripping system while the load was measured by a load cell that was connected to the end of one gripping jaw. Displacement was obtained using a Linear Variable Differential Transformer (LVDT) which was attached to the two gripping jaws to measure the change of gauge length.
Before testing, the thickness and width of a ribbon were carefully measured for at least three times at different locations in the gauge length. The average values were then recorded as gauge thickness and width, and used as input parameters for subsequent stress and strain calculation. The initial gauge length for tensile testing was set at ˜2.50 mm with the exact value determined after the ribbon was fixed, by accurately measuring the ribbon span between the front faces of the two gripping jaws. All tests were performed under displacement control, with a strain rate of ˜0.001 s−1.
In Table 7, a summary of the tensile test results including total elongation, yield strength, ultimate tensile strength, Young's Modulus, Modulus of Resilience, and Modulus of Toughness are shown for each alloy of Table 1 when melt-spun at both 16 and 10.5 m/s. Note that each distinct sample was measured in triplicate since occasional macrodefects arising from the melt-spinning process can lead to localized areas with reduced properties. The results shown in Table 7 have not been adjusted for machine compliance.
For the tensile measurements shown in Table 7, the data can be corrected to adjust for machine compliance coefficient and deviations in cross sectional area from rectangular cross sections. The corrected data which represents the most accurate tensile results are shown in Table 8. As can be seen the tensile strength values are relatively high and vary from 0.36 to 2.77 GPa while the total elongation values are also very significant for reduced length scale microstructures and vary from 0.65 to 4.61%.
Proposed Mechanism
The following mechanism for microstructural formation has been developed to qualify the current results including the measured high elongation and the four distinct types of bending behavior observed in the melt-spun alloys. Note that these models are developed to coordinate the results but in no way are construed to limit the features of specific details of potentially more complex interactions. Additionally, the mechanism of microstructural formation and specific structural features may be relevant to a wide variety of metallic glass chemistries made with different base metals such as nickel, cobalt, magnesium, titanium, molybdenum, rare earths, etc.
If nucleation is completely avoided during solidification, a metallic glass structure may be formed. The metallic glass structure at room temperature is known to deform upon the application of a tensile stress by a localized inhomogeneous mechanism called shear banding resulting in brittle failure. Current research shows, high elongation and high bending strains occur only in specific samples which have significant and measurable amounts of metallic glass present. However, the presence of metallic glass alone is not expected nor believed to be the source of high elongation. Based on current results, it is believed that crystalline phase formation during solidification may occur in two distinct modes, Glass devitrification and Spinodal decomposition. Glass Devitrification may be understood to occur through nucleation and growth resulting from a high driving force in the supercooled melt which leads to a high nucleation frequency, limited time for growth and the achievement of nanoscale phases. Depending on the specific cooling rate, the devitrification transformation can occur completely (for Example see
For the studied alloys, it is believed that examples of spinodal decomposition in various forms were shown including microconstituent bands (for example see
In Table 6, for the alloy studied, the 180° bend tests were correlated and, as stated earlier, revealed 4 distinct types of behavior which were; Type 1 Behavior: Not bendable flat in either direction, Type 2 Behavior: Bendable flat in one direction with wheel side out, Type 3 Behavior: Bendable flat in one direction with free side out, and Type 4 Behavior: Bendable flat in both directions. The bending behavior illustrates material response over a fairly large area of bending and along the length of the ribbon since the bending response generally occurs along the entire length of the ribbon with the exception of isolated spots which, in most cases, can be attributed to macrodefects arising from the melt-spinning process. Note that during 180° bending the outside of the ribbon is placed into tension while the inside of the ribbon is placed into compression. While metallic glasses and other brittle structures may perform well in compression, in tension where shear bands (i.e., in situ tensile deformations) and cracks can propagate, metallic glasses may fail in a brittle manner. Thus, the different bending results may indicate differences in structure between the free, center, and wheel sides of the ribbons. The X-ray results shown in
Elongation of >0.65% is expected to be achieved through the interaction of the shear bands formed in the glass matrix with various crystalline features. While all crystalline features may be expected to provide some pinning or interaction with the domain walls based on the entirety of the results, it is believed that the most effective pinning/blunting, and shearing is occurring from the spinodal microconstituent regions. Thus, the following models are proposed to explain observed behavior. Note that the cooling rate at the wheel surface is the fastest due to conductive heat transfer to the copper wheel, followed by the free surface due to conductive/radiative heat transfer to the helium gas, and then followed by the center of the ribbon which is limited by thermal conductivity to the outside surfaces.
Type 1 Behavior Model
In
Type 2 Behavior Model
In
Type 3 Behavior Model
In
Type 4 Behavior Model
In
Using high purity elements, fifteen gram charges of the PC7E4C3 chemistry were weighed out according to the atomic ratio's in Table 1. The mixture of elements was placed onto a copper hearth and arc-melted into ingots using ultrahigh purity argon as a cover gas. After mixing, the resulting ingots were cast into a figure shape appropriate for melt-spinning. The cast fingers of PC7E4C3 were then placed into a quartz crucible with a hole diameter nominally at 0.81 mm. The ingots were heated up by RF induction and then ejected onto a rapidly moving 245 mm copper wheel traveling at wheel tangential velocities of 16 and 10.5 m/s. To further examine the ribbon structure, scanning electron microscopy (SEM) was done on selected PC7E4C3 ribbon samples. Melt spun ribbons were mounted in a standard metallographic mount with several ribbons held using a metallography binder clip. The binder clip containing the ribbons was set into a mold and an epoxy is poured in and allowed to harden. The resulting metallographic mount was ground and polished using appropriate media following standard metallographic practices. The structure of the samples was observed using a Zeiss EVO-60 scanning electron microscope with an electron beam energy of 17.5 kV, a filament current of 2.4 A, a spot size setting of 800. As shown in
Using high purity elements, a fifteen gram charge of the PC7E7 alloy was weighed out according to the atomic ratio's in Table 1. The mixture of elements was placed onto a copper hearth and arc-melted into an ingot using ultrahigh purity argon as a cover gas. After mixing, the resulting ingot was cast into a finger shape appropriate for melt-spinning. The cast fingers of PC7E7 were then placed into a quartz crucible with a hole diameter nominally at 0.81 mm. The ingots were heated up by RF induction and then ejected onto a rapidly moving 245 mm copper wheel traveling at a wheel tangential velocities of 16 m/s. The ribbon was cut into pieces and then tested in tension and the resulting tensile test stress/strain data from one test is shown in
Using high purity elements, a fifteen gram charge of the PC7E7 alloy was weighed out according to the atomic ratio's in Table 1. The mixture of elements was placed onto a copper hearth and arc-melted into an ingot using ultrahigh purity argon as a cover gas. After mixing, the resulting ingot was cast into a finger shape appropriate for melt-spinning. The cast fingers of PC7E7 were then placed into a quartz crucible with a hole diameter nominally at 0.81 mm. The ingots were heated up by RF induction and then ejected onto a rapidly moving 245 mm copper wheel traveling at a wheel tangential velocities of 10.5 m/s. A piece of typical ribbon was then selected for TEM and was cut into three consecutive short segments. For each segment, the ribbons were mechanically thinned from their original thickness to approximately 10 microns using fine-grit sandpaper followed by polishing using 5 micron and 0.3 micron alumina powder on felt pads with water as a lubricant. The thinning of the three samples is shown in
Using high purity elements, a fifteen gram charge of the PC7E7 alloy was weighed out according to the atomic ratio's in Table 1. The mixture of elements was placed onto a copper hearth and arc-melted into an ingot using ultrahigh purity argon as a cover gas. After mixing, the resulting ingot was cast into a finger shape appropriate for melt-spinning. The cast fingers of PC7E7 were then placed into a quartz crucible with a hole diameter nominally at 0.81 mm. The ingots were heated up by RF induction and then ejected onto a rapidly moving 245 mm copper wheel traveling at a wheel tangential velocity of 10.5 m/s. Sample of ribbon were then etched with a 2% bromine water solution. The structure of the etched sample was observed using an EVO-60 scanning electron microscope manufactured by Carl Zeiss SMT Inc. Typical operating conditions were electron beam energy of 17.5 kV, filament current of 2.4 A, and spot size setting of 800. In
Using high purity elements, 15 g alloy feedstocks of the targeted alloys were weighed out according to the atomic ratio's provided in Table 9. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and remelted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. The resulting fingers were then placed in a melt-spinning chamber in a quartz crucible with a hole diameter of ˜0.81 mm. The ingots were melted in a ⅓ atm helium atmosphere using RF induction and then ejected onto a 245 mm diameter copper wheel which was traveling at tangential velocities of 10.5 m/s. Bending testing (180° of the as-spun ribbon samples was done on each sample and the results were correlated in Table 10. As shown, depending on the alloy when processed on the particular conditions listed, the bending response was found to vary, with four types of behavior observed; Type 1 Behavior—not bendable without breaking, Type 2 Behavior—bendable on one side with wheel side out, Type 3 Behavior—bendable on one side with free side out, and Type 4 Behavior—bendable on both sides. In Table 11, a summary of the tensile test results including total elongation, yield strength, ultimate tensile strength, Young's Modulus, Modulus of Resilience, and Modulus of Toughness are shown for each alloy of Table 8 when melt-spun at 10.5 m/s. Note that each distinct sample was measured in triplicate since occasional macrodefects arising from the melt-spinning process can lead to localized stresses reducing properties. The results shown in Table 11 have not been adjusted for machine compliance.
For the tensile measurements shown in Table 11, the data can be corrected to adjust for machine compliance coefficient and deviations in cross sectional area from rectangular cross sections. The corrected data which represents the most accurate tensile results are shown in Table 12. As can be seen the tensile strength values are high and vary from 0.40 to 3.47 GPa while the total elongation values are very significant for reduced length scale microstructures and vary from 0.65 to 7.06%.
Using commercial purity feedstock including ferroadditives, 15 g alloy feedstocks of the targeted alloys were weighed out according to the atomic ratio's provided in Table 13. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and remelted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. For each ingot, the density was measured using the Archimedes principle and the results are shown in Table 14. As shown, the densities were found to vary from 7.28 to 7.81 g/cm3. The resulting arc-melted ingots were then placed in a melt-spinning chamber in a quartz crucible with a hole diameter of ˜0.81 mm. The ingots were melted in a air using RF induction and then ejected with a melt superheat of 150° C. and a chamber pressure of 280 mbar onto a 245 mm diameter copper wheel which was traveling at tangential velocities of 25 m/s. Long ribbon lengths typically from 0.7 to 1.5 mm in width were obtained. The thickness of the ribbons produced was then measured in a micrometer and the results are tabulated in Table 14. As shown, the thickness was dependant on alloy chemistry and was found to vary from 37 to 55 μm. Bending testing (180°) of the as-spun ribbon samples were done on each sample and the results were correlated in Table 9. As shown, depending on the alloy when processed on the particular conditions listed, the bending response was found to vary but the primary response was Type 4 Behavior (i.e. bendable on both sides).
Thermal analysis was done on the as-solidified ribbons of Table 13 using a Perkin Elmer DTA-7 system with the DSC-7 option. Differential thermal analysis (DTA) and differential scanning calorimetry (DSC) was performed at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultrahigh purity argon. In Table 15, the DSC data related to the glass to crystalline transformation is shown for the alloys that have been melt-spun in air at 25 m/s. All of the samples were found to contain a significant fraction of glass. The glass to crystalline transformation occurs in either one stage or two stages in the range of temperature from 452 to 595° C. and with enthalpies of transformation from −22.8 to 115.8 J/g.
In Table 16, a summary of the tensile test results including gage dimensions, elongation, yield breaking load, strength and Young's Modulus are shown for each alloy of Table 13. Note that each distinct sample was measured in triplicate since occasional macrodefects arising from the melt-spinning process can lead to localized stresses reducing properties. As can be seen the total elongation values are significant and vary from 1.97 to 4.78% with high tensile strength values from to GPa. Young's Modulus was found to vary from 1.12 to 2.92 GPa. Note that the results shown in Table 16 have been adjusted for machine compliance and geometric cross sectional area.
Using commercial purity feedstock including ferroadditives, 15 g alloy feedstocks of the C01F03 and C01B03 alloys were weighed out according to the atomic ratio's provided in Table 13. The feedstock material was then placed into the copper hearth of an arc-melting system. The feedstock was arc-melted into an ingot using high purity argon as a shielding gas. The ingots were flipped several times and remelted to ensure homogeneity. After mixing, the ingots were then cast in the form of a finger approximately 12 mm wide by 30 mm long and 8 mm thick. To show variability in properties, the alloys were process into ribbons using different parameters as shown in Table 17. Note that the C01F03 alloy was processed using the MS45 and MS58 parameter while the C01B03 alloy was processed using the MS45, MS50, and MS55 parameter. Thermal analysis was done on the as-solidified ribbons of Table 13 using a Perkin Elmer DTA-7 system with the DSC-7 option at a heating rate of 10° C./minute with samples protected from oxidation through the use of flowing ultrahigh purity argon. In Table 19, the results of the DSC analysis is shown. As indicated, the onset temperatures, peak temperatures, and enthalpies can vary with process parameter.
In Table 20, a summary of the tensile test results including gage dimensions, elongation, yield breaking load, strength and Young's Modulus are shown for each alloy of Table 13. Note that each distinct sample was measured in triplicate since occasional macrodefects arising from the melt-spinning process can lead to localized stresses reducing properties. As can be seen the tensile properties can vary dramatically as a function of processing parameter. Note that the results shown in Table 16 have been adjusted for machine compliance and geometric cross sectional area.
The results of this case example, clearly show some of the variability with respect to physical property changes of the alloys with respect to process parameters. In an illustrated case for example, the C01B03 alloy went from a ductile sample when processed at the MS50 parameter to a brittle sample when processed using the MS55 parameter. This change is consistent with the proposed mechanism which shows that only with specific structures does a ductile sample result. It is believed that a much broader range in properties could be obtained with the identified alloys by optimizing process parameters further. Additionally, this shows that additionally chemistry variations are possible by commensurately changing process parameters. Note that this change in structure and properties through processing alterations is a well established basis in modern metallurgy.
The foregoing description of several methods and embodiments has been presented for purposes of illustration. It is not intended to be exhaustive or to limit the claims to the precise steps and/or forms disclosed, and obviously many modifications and variations are possible in light of the above teaching. It is intended that the scope of the invention be defined by the claims appended hereto.
This application claims the benefit of U.S. Provisional Patent Application Ser. No. 61/107,037, filed on Oct. 21, 2008, which is fully incorporated herein by reference.
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| Number | Date | Country | |
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| 20100154942 A1 | Jun 2010 | US |
| Number | Date | Country | |
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| 61107037 | Oct 2008 | US |
