Low Thermal Expansion High-Strength Alloy and Method of manufacturing the same

Abstract

A low thermal expansion, high-strength alloy is disclosed that possesses a coefficient of thermal expansion similar to that of conventional Invar alloys while having superior mechanical properties. The grain shape is controlled through rolling, and the size and shape of the grains in the matrix are controlled during the heat treatment process. Furthermore, the coefficient of thermal expansion and mechanical strength are controlled by forming precipitates in the matrix phase or at the grain boundaries depending on the heat treatment temperature.

Description

Technology Field

[0001] The present invention relates to an alloy that achieves high strength while having a low coefficient of thermal expansion, and more specifically, to a low thermal expansion high-strength alloy that has higher strength than conventional alloys and allows for the control of the coefficient of thermal expansion by controlling the composition of the alloy's constituent metals, and a method for manufacturing the same. Background Technology

[0002] Low thermal expansion alloys possess the characteristic of having a low coefficient of thermal expansion. Since low thermal expansion alloys exhibit minimal dimensional change due to temperature variations, they offer the advantage of maintaining shape retention and precision in precision environments or devices.

[0003] Since metal materials have the characteristic of expanding or contracting when temperature changes, errors caused by thermal expansion become a significant problem in precision instruments or measuring devices where minute dimensional errors must be controlled. Furthermore, in environments where different materials such as metals and ceramics are assembled, differences in thermal expansion coefficients can lead to cracking and detachment of the materials due to thermal shock.

[0004] To address this, alloys with low thermal expansion characteristics are used, and a representative alloy is Invar Alloy. Invar Alloy is an alloy composed mainly of iron and nickel, and it has the advantage of a very low coefficient of thermal expansion with temperature changes. For example, it exhibits a coefficient of thermal expansion of 1 ppm / °C over a temperature change from room temperature to 200°C. However, low mechanical strength is pointed out as a weakness of Invar Alloy. Although there are differences depending on manufacturing methods such as annealing and cold working, it exhibits a maximum tensile strength of approximately 430 MPa and a yield strength of approximately 240 MPa, which are lower than those of ordinary steel.

[0005] Therefore, conventional low thermal expansion alloys such as Invar alloys are primarily used in areas requiring precision and thermal stability rather than in areas requiring structural load endurance.

[0006] However, applications requiring high mechanical strength along with thermal stability are on the rise. For example, the aerospace sector demands alloys with robust physical properties that withstand mechanical vibrations and shocks even in environments with rapid temperature changes at high altitudes. Furthermore, precision optical systems, semiconductor lithography equipment, and power plant facilities require low-thermal-expansion, high-strength alloys that suppress thermal expansion while resisting mechanical wear and vibration.

[0007] To satisfy the required material properties, alloys such as Super Invar or Kovar, which feature low thermal expansion characteristics and reinforced strength, are being introduced. However, as these are alloys designed with the primary focus on controlling thermal expansion, they possess lower mechanical strength compared to general structural steel.

[0008] Therefore, new alloys that possess high mechanical strength while allowing control of the coefficient of thermal expansion through compositional adjustment are still in demand. The problem to be solved

[0009] The first technical objective of the present invention is to provide a low thermal expansion, high-strength alloy that has higher strength than conventional Invar alloys while controlling the coefficient of thermal expansion.

[0010] In addition, the second technical objective of the present invention is to provide a method for manufacturing a low thermal expansion, high-strength alloy to achieve the first technical objective. means of solving the problem

[0011] The present invention, for achieving the first technical objective described above, provides a low thermal expansion high-strength alloy characterized by comprising Fe: 43.2 at.% to 62.5 at.%, Co: 6.8 at.% to 33.4 at.%, Ni: 10.3 at.% to 29.3 at.%, Cr: 0.5 at.% to 3.8 at.%, Mo: 0.9 at.% to 9.7 at.%, and other unavoidable impurities, having a face-centered cubic matrix structure through annealing, and containing σ phase or μ phase precipitates within the grain boundaries.

[0012] The present invention, for achieving the second technical objective described above, provides a method for manufacturing a low thermal expansion high-strength alloy comprising the steps of: performing a homogenization heat treatment on a casting in which a metal raw material is melted and cast; performing cold rolling after the homogenization heat treatment; and controlling crystal grains and precipitates forming the matrix of the alloy through annealing of the cold-rolled alloy. Effects of the invention

[0013] According to the present invention described above, strength enhancement is maximized through a multi-component alloy, and the coefficient of thermal expansion exhibits characteristics similar to those of conventional Invar-based alloys. To enhance strength, fine nano-sized precipitates are formed within a face-centered cubic matrix. These precipitates are formed at grain boundaries or within grains through homogenization heat treatment and annealing processes; while their size or shape may vary depending on the heat treatment conditions, they possess a fine size that is not coarsenic. Depending on the heat treatment conditions, partial recrystallization may occur, resulting in the formation of a heterostructure microstructure. In other words, a combination of the heterostructure microstructure and nano-precipitates ensures high strength and excellent ductility.

[0014] In addition, the formed face-centered cubic matrix exhibits low thermal expansion characteristics due to thermal expansion caused by lattice vibrations and lattice shrinkage effects caused by spontaneous volume magnetostriction through the multi-component compositional design of the constituent elements, and the fine-sized precipitates generated internally do not significantly reduce the coefficient of thermal expansion of the alloy. Therefore, a coefficient of thermal expansion similar to that of existing Invar-based alloys can be obtained. Brief explanation of the drawing

[0015] FIG. 1 is a flowchart illustrating a method for manufacturing a low thermal expansion high-strength alloy according to a preferred embodiment of the present invention. Figure 2 is a graph showing the change in hardness according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention. Figure 3 is a graph showing the change in hardness according to the annealing time of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention. Figure 4 is an X-ray diffraction (XRD) graph according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention. Figure 5 shows field emission-scanning electron microscope (FE-SEM) backscattered electron (BSE) images of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention as a function of annealing temperature. Figure 6 is an XRD graph analyzing the crystal structure of precipitates in a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention. Figure 7 shows the electron backscatter diffraction (EBSD) map of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention. FIG. 8 is a graph showing the results of a tensile test according to Measurement Example 1 of the present invention. FIG. 9 is a graph showing the magnetic properties of alloy specimens according to Measurement Example 1 of the present invention. FIG. 10 is a graph showing the coefficient of thermal expansion of alloy specimens according to Measurement Example 1 of the present invention. Figure 11 is a graph showing the change in hardness according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention. FIG. 12 is a graph showing the change in hardness according to the annealing time of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention. Figure 13 shows FE-SEM BSE images according to the annealing temperature of the low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention. Figure 14 is an XRD graph analyzing the crystal structure of precipitates in a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention. Figure 15 shows the EBSD map of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention. FIG. 16 is a graph showing the results of a tensile test according to Measurement Example 2 of the present invention. FIG. 17 is a graph showing the magnetic properties of alloy specimens according to Measurement Example 2 of the present invention. FIG. 18 is a graph showing the coefficient of thermal expansion of alloy specimens according to Measurement Example 2 of the present invention. FIG. 19 is a graph showing the hardness of alloys according to Manufacturing Examples 3 and 4 of the present invention. FIG. 20 is a graph comparing the hardness of the commercial Invar 36 alloy of the present invention and the alloys of Manufacturing Examples 3 and 4. Figure 21 is an XRD graph analyzing the crystal structures of Manufacturing Examples 3 and 4 of the present invention. Figure 22 shows FE-SEM BSE images of the alloys of Manufacturing Examples 3 and 4 of the present invention according to the annealing temperature. FIG. 23 is a graph showing the tensile test results of alloy specimens provided in Manufacturing Examples 3 and 4 according to Measurement Example 3 of the present invention. FIG. 24 is a graph showing the coefficients of thermal expansion of alloy specimens of Manufacturing Examples 3 and 4 measured according to Measurement Example 3 of the present invention. FIG. 25 is a graph showing the hardness according to the annealing temperature of alloys according to Manufacturing Example 1 and Manufacturing Example 5 of the present invention. FIG. 26 is a graph showing the hardness according to the annealing time of alloys according to Manufacturing Example 1 and Manufacturing Example 5 of the present invention. FIG. 27 is an XRD graph analyzing the crystal structure of alloys according to Manufacturing Example 5 of the present invention. FIG. 28 is FE-SEM BSE images of alloys according to Manufacturing Example 5 of the present invention. FIG. 29 is an energy dispersive X-ray spectroscopy (EDX) graph analyzing the components of the precipitate according to Preparation Example 5 of the present invention. FIG. 30 is an EBSD Inverse Pole Figure (IPF) map and Inverse Pole Figures of the Rolling Direction (RD) of the alloys according to Manufacturing Example 5 of the present invention. FIG. 31 is a graph showing the results of a tensile test at room temperature according to Measurement Example 4 of the present invention. FIG. 32 is -196 according to Measurement Example 4 of the present invention This is a graph showing the results of the tensile test in. FIG. 33 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 1 according to a simulation example of the present invention. FIG. 34 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 1 according to a simulation example of the present invention. FIG. 35 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 2 according to a simulation example of the present invention. FIG. 36 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 2 according to a simulation example of the present invention. FIG. 37 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 3 according to the simulation example of the present invention. FIG. 38 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 3 according to a simulation example of the present invention. FIG. 39 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 4 according to the simulation example of the present invention. FIG. 40 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 5 according to the simulation example of the present invention. FIG. 41 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 5 according to a simulation example of the present invention. FIG. 42 shows Fe according to a composition design example of the present invention. 42.91 Ni 10.3 Co 31 Cr 3.8 Mo 9.7 C 0.29 This is a graph predicting the phase change of the alloy. FIG. 43 shows Fe according to a composition design example of the present invention 39.81 Ni 9.73 Co 32.93 Cr 2.33 Mo 13.00 V 2.0 C 0.2 This is a graph predicting the phase change of the alloy. Specific details for implementing the invention

[0016] The present invention is susceptible to various modifications and may take various forms; therefore, specific embodiments are illustrated in the drawings and described in detail in the text. However, this is not intended to limit the invention to the specific disclosed forms, and it should be understood that the invention includes all modifications, equivalents, and substitutions that fall within the spirit and scope of the invention. Similar reference numerals have been used for similar components in the description of each drawing.

[0017] Unless otherwise defined, all terms used herein, including technical or scientific terms, have the same meaning as generally understood by those skilled in the art to which the present invention pertains. Terms such as those defined in commonly used dictionaries should be interpreted as having a meaning consistent with their meaning in the context of the relevant technology, and should not be interpreted in an ideal or overly formal sense unless explicitly defined in this application.

[0018] Hereinafter, preferred embodiments of the present invention will be described in more detail with reference to the attached drawings.

[0020] Examples

[0021] The low thermal expansion, high strength alloy of this embodiment is Fe a Co b Ni c Cr d Mo e has.

[0022] In the above low thermal expansion high strength alloy, the sum of the at.% of the elements, a+b+c+d+e, is 100 at% including other unavoidable impurities.

[0023] The content of the elements constituting the alloy is as follows.

[0024] Fe : 43.2 at.% to 62.5 at.%,

[0025] Co : 6.8 at.% to 33.4 at.%,

[0026] Ni: 10.3 at.% to 29.3 at.%,

[0027] Cr : 0.5 at.% to 3.8 at.%,

[0028] It has an alloy composition containing Mo: 0.9 at.% to 9.7 at.% and other unavoidable impurities.

[0029] In addition, the low thermal expansion, high strength alloy of this embodiment is Fe a Co b Ni c Cr d Mo e V f has.

[0030] In the above low thermal expansion high strength alloy, the sum of the at.% of the elements a+b+c+d+e+f is 100 at% including other unavoidable impurities.

[0031] The content of the elements constituting the alloy is as follows.

[0032] Fe : 43.2 at.% to 62.5 at.%,

[0033] Co : 6.8 at.% to 33.4 at.%,

[0034] Ni: 10.3 at.% to 29.3 at.%,

[0035] Cr : 0.5 at.% to 3.8 at.%,

[0036] Mo : 0.9 at.% to 9.7 at.%

[0037] V: It has an alloy composition containing 0 at% to 2.0 at% and other unavoidable impurities.

[0038] In addition to the composition described above, the alloy may further contain C (carbon), and said carbon may be added in the alloy in an amount of 0 at.% to 0.29 at.%. When carbon is added, Fe a Co b Ni c Cr d Moe The combined content of other metals, carbon, and other unavoidable impurities in the alloy is 100 at.%.

[0039] Below, the numerical values ​​of the elements constituting the alloy, expressed in at.% units, are represented as the composition, content, or fraction of the alloy.

[0040] Fe is the most abundant base element in the alloy and forms a face-centered cubic matrix within the aforementioned compositional range. In this specification, "matrix" refers to the base of the alloy and is used interchangeably with "base" or "matrix" below. Fe has excellent mechanical properties, the ability to undergo phase transformation into various crystal structures depending on annealing conditions and phase stability, and the advantage of excellent miscibility with other metallic elements.

[0041] Co, along with Ni, which will be discussed later, is an important element that determines the coefficient of thermal expansion. In particular, it is an indispensable element for obtaining a low coefficient of thermal expansion. If the content of Co exceeds 33.4 at.%, the coefficient of thermal expansion in the alloy increases rapidly, and if it is less than 6.8 at.%, changes or increases in the coefficient of thermal expansion occur depending on the proportion of other metal components, making it difficult to control a stable coefficient of thermal expansion.

[0042] Ni is an important element that, along with Co, determines the coefficient of thermal expansion in the alloy. Although it exhibits the lowest coefficient of thermal expansion at 36 at.%, this can reduce the strengthening effect as it decreases the composition of other elements that contribute to strength improvement. Therefore, to simultaneously achieve a low coefficient of thermal expansion and high mechanical strength, the Ni content is controlled within an optimal range. If the Ni content is less than 10.3 at.%, the coefficient of thermal expansion increases, and if it exceeds 29.3 at.%, the sensitivity of the alloy decreases rapidly.

[0043] Cr is an element used to improve the corrosion resistance and strength of the alloy. In this embodiment, if the Cr content is less than 0.5 at.%, the strengthening effect is small, and if it exceeds 3.8 at.%, precipitates are excessively formed, increasing brittleness and lowering thermal expansion properties.

[0044] Mo is an element added to form precipitates, thereby improving strength. Additionally, Mo remaining in the matrix exhibits a solid solution strengthening effect. As the Mo content increases, the number of precipitates increases, hindering grain growth and suppressing recrystallization. If the Mo content exceeds 9.7 at.%, the ductility of the alloy decreases, and the coefficient of thermal expansion increases. Furthermore, if the Mo content is less than 0.9 at.%, precipitates are not formed, and the solid solution strengthening effect is reduced.

[0045] V promotes the formation of precipitates and ensures that the precipitates are evenly dispersed within the alloy. This can lead to an improvement in the strength of the alloy. Additionally, it promotes the formation of Mo-rich precipitates, thereby inducing a solid solution strengthening effect. However, if the content exceeds 2.0 at.%, there is a problem where an excessive amount of carbides is formed, leading to increased brittleness. However, in this embodiment, V can be applied selectively within the alloy.

[0046] In addition, C can also be selectively applied within the alloy and may contribute to the improvement of the alloy's strength by additionally forming carbide precipitates. If the C content exceeds 0.29 at.%, the coefficient of thermal expansion of the alloy may increase. That is, the C content may be included in the alloy in the range of 0 to 0.29 at.%, where the sum of the at.% values ​​of other metal elements constituting the alloy, other unavoidable impurities, and the at.% of C is 100.

[0047] FIG. 1 is a flowchart illustrating a method for manufacturing a low thermal expansion high-strength alloy according to a preferred embodiment of the present invention.

[0048] Referring to FIG. 1, homogenization annealing is performed on an alloy of the composition described above (S100).

[0049] Before homogenization annealing is performed, the raw materials constituting the alloy are placed in a crucible and melted and cast through heating. Homogenization annealing is an annealing process that eliminates non-uniformity of components formed inside the alloy after melting and makes the metal structure uniform. In addition, the initial grain size of the alloy can be controlled by adjusting the annealing conditions.

[0050] The above homogenization annealing is carried out at a temperature lower than the melting point of the alloy and serves to reduce the non-uniformity of components formed during casting.

[0051] For example, the homogenization annealing described above may be carried out at 1200°C for 6 hours, and the homogenization annealing conditions may vary depending on the composition of the alloy, added elements, the shape and weight of the casting, the cooling rate, etc.

[0052] For the alloy that has undergone homogenization annealing, cold rolling is performed at a reduction rate within a certain range (S200). Even if the reduction rate and rolling conditions are different (e.g., hot rolling), a heteromicrostructure in which the core precipitates of this development are induced can be obtained.

[0053] The above cold rolling is performed below the recrystallization temperature of the alloy and, preferably, can be performed at room temperature. The rolling process of the present invention can be extended to the form of cold rolling or hot rolling, but for convenience of explanation, the following examples are described based on room temperature cold rolling conditions. Grains that have grown or coarseened through homogenization annealing are plastically deformed through cold rolling, thereby changing their structure and crystallographic orientation, and through this, the strength of the alloy is improved.

[0054] The above cold rolling is performed to refine the grain size of the alloy or control precipitates, and the reduction rate can be set to various values ​​depending on the alloy composition and process conditions. For example, in this embodiment, a reduction rate of approximately 50% was applied. Subsequently, annealing is performed on the cold-rolled alloy (S300).

[0055] Depending on the annealing temperature, the size and shape of the crystal grains can be controlled, and the crystal structure, shape, size, and fraction of the precipitates can also be controlled. The precipitates are formed when supersaturated elements in a solid solution state are separated into fine particles during the annealing process and have a chemical composition or crystal structure different from that of the substrate or matrix.

[0056] The above annealing can remove residual stress within the rolled alloy and induce recrystallization based on the dislocation density accumulated within it.

[0057] The annealing temperature can be performed at 500℃ or higher, and the balance between mechanical strength and ductility must be considered. At this time, treatment in the range of 500℃ to 950℃ is preferable. However, annealing at higher temperatures may reduce mechanical strength due to grain growth and the dissolution of precipitates, but it can contribute to lowering the coefficient of thermal expansion. The alloy after annealing treatment can be cooled by various methods such as water quenching, air cooling, or furnace cooling, and among these, an alloy was obtained that simultaneously secures a low coefficient of thermal expansion and high strength by applying water quenching.

[0059] Preparation Example 1: Fe 43.2 Ni 10.3 Co 33.0 Cr 3.8 Mo 9.7 manufacturing

[0060] Metal raw materials of Fe, Ni, Co, Cr, and Mo with a purity of 99% or higher are prepared. The mixing ratio of these metals is 43.2 : 10.3 : 33.0 : 3.8 : 9.7 in at.% units.

[0061] After placing the metal raw material into a ceramic crucible, it is heated to 1300°C to 1400°C for 4 to 15 minutes using induction heating to melt it, and then cast into a carbon mold. However, those skilled in the art may change the heating temperature and time depending on the capacity of the ceramic crucible, etc.

[0062] The casting is homogenized and heat-treated at 1200°C for 6 hours. Then, the homogenized alloy is cold-rolled with a reduction rate of 50%.

[0063] The above rolled alloys are annealed at 550°C, 650°C, 750°C, 850°C, 950°C, and 1050°C, respectively. The annealing time is set to 1 hour. Afterward, water quenching is performed to obtain a low thermal expansion high-strength alloy.

[0064] Figure 2 is a graph showing the change in hardness according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0065] Referring to Fig. 2, the sample annealed at 550°C exhibits the highest hardness. Additionally, it is observed that the hardness decreases as the annealing temperature increases. In the annealing temperature range above 550°C, as the temperature rises, recrystallization due to annealing proceeds, and the hardness of the alloy decreases.

[0066] Subsequently, to determine the tendency of precipitates within the alloy, annealing conditions are set to 550°C, 750°C, and 950°C. Then, to check the change in hardness over time under the set annealing conditions, annealing is performed at different times.

[0067] Figure 3 is a graph showing the change in hardness according to the annealing time of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0068] Referring to Fig. 3, the alloy exhibits the highest hardness when the annealing time is 1 hour. If the annealing time exceeds 1 hour, a decrease in hardness is observed. This is interpreted as precipitation within the alloy occurring within the first hour of annealing, and subsequent annealing periods causing the precipitates to coarsen, thereby further lowering the hardness of the alloy.

[0069] However, the crystal structure of the matrix needs to be verified at the set temperature regardless of the annealing time.

[0070] Figure 4 is an XRD graph according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0071] Referring to Fig. 4, a face-centered cubic crystal structure is confirmed at all three selected annealing temperatures regardless of time. In particular, a strong peak of the face-centered cubic structure is formed in the sample annealed at 550°C. This indicates that the crystal structure of the matrix at 550°C is well formed as a single phase of the face-centered cubic structure.

[0072] Figure 5 shows FE-SEM BSE images according to the annealing temperature of the low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0073] Through the above images, the phases appear with different brightness, through which the particle distribution or phase formation of the alloy can be identified.

[0074] Referring to FIG. 5, images (a) and (b) show a sample annealed at 550°C for 1 hour, images (c) and (d) show a sample annealed at 750°C for 1 hour, and images (e) and (f) show a sample annealed at 950°C for 1 hour.

[0075] In images (a) and (b), fine σ (sigma) phases are observed within the matrix. In images (a) and (b), they appear as irregular particles, and their size is so fine that they are barely observable in high-magnification scanning electron microscope images. The σ phase is characterized by precipitation mainly along grain boundaries. However, the formation of the σ phase does not occur in large quantities, and only fine formations are observed.

[0076] Referring to images (c) and (d), the σ phase grows and appears in a plate-like form, and the μ (mu) phase appears mixed within the σ phase. The μ phase is interpreted as a phase that appears upon the introduction of Mo into the alloy.

[0077] Referring to images (e) and (f), the grains recrystallize, and multiple μ phases precipitate around the grain boundaries. Additionally, smaller μ phases precipitate within the grains.

[0078] The image in Fig. 5 above has the same trend as the hardness graphs in Figs. 2 and 3. That is, when a large amount of σ phase and μ phase precipitates in the grains or grain boundaries, the hardness of the alloy decreases.

[0079] Figure 6 is an XRD graph analyzing the crystal structure of precipitates in a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0080] Referring to FIG. 6, graph (a) is an XRD graph of an alloy that has been annealed at 550°C for 1 hour. In the graph, the size of the σ phase precipitates is so fine that they are not detected as XRD phases.

[0081] In graph (b), the σ phase of the tetragonal structure is clearly observed.

[0082] In addition, a peak of the μ phase with a rhombohedral structure is observed in graph (c). Topologically, it is a close-packed phase, and both the μ phase and the sigma phase are included in it.

[0083] The above results are consistent with the image in Fig. 5, and it is confirmed that the hardness of the alloy decreases at high annealing temperatures.

[0084] Table 1 below shows the compositional analysis of the alloy matrix and precipitates via Point-EDS.

[0085] annealing temperature type Fe Co Ni Cr Mo 550℃ matrix 43.1±0.2 34.5±0.3 10.0±0 3.5±0 8.9±0.1 precipitates 41.3±0.5 33.7±0.1 10.7±0.3 3.9±0.1 10.3±0 750℃ matrix 43.1±0.1 34.4±0.3 10.1±0.2 3.4±0.1 9.0±0.1 precipitates 42.2±0.4 34.1±0.3 9.7±0.1 3.6±0.1 10.4±0.7 950℃ matrix 43.9±0.6 34.9±0.4 10.3±0.2 3.5±0.1 7.4±0.7 precipitates 35.0±3.0 31.0±1.3 6.4±1.2 3.4±0.1 24.2±5.6

[0086] In Table 1 above, the unit of composition is at.%, and these are alloys annealed for 1 hour. The metallic elements whose content ratio in the precipitates relative to the matrix increased at 550°C and 750°C are Cr and Mo. That is, Cr and Mo appear abundant in the precipitates. This is because the precipitates are in the σ phase, and the σ phase is based on intermetallic compounds rich in Cr and Mo.

[0087] In addition, at an annealing temperature of 950°C, the proportion of Mo element relative to the matrix increased most significantly in the precipitates. Since the μ phase is a crystalline phase due to the high content of Mo, it is confirmed that a large amount of μ phase precipitates were formed in the alloy annealed at 950°C.

[0088] Figure 7 shows the EBSD map of a low thermal expansion high-strength alloy according to Manufacturing Example 1 of the present invention.

[0089] Image (a) is an Inverse Pole Figure (IPF) map, which allows verification of the direction of the crystal axes. Image (b) is a Misorientation Angle map, which confirms the difference in azimuth angles between grains. Image (c) is a Kernel Average Misorientation (KAM) map, which identifies the deformation or stress state of the grains. Additionally, Image (d) is a Grain Orientation Spread (GOS) map, which identifies the consistency of orientation within the grains.

[0090] Referring to the images in Fig. 7 above, the alloy annealed at 550°C and 750°C does not undergo recrystallization, so large, plastically deformed grains are observed with the rolled microstructure generated by cold rolling remaining.

[0091] However, in the alloy annealed at 950°C, partial recrystallization proceeded, and a bimodal structure in which fine and coarse grains coexist was confirmed. However, it can be said that recrystallization under these heat treatment conditions occurred only to a limited extent due to the pinning effect caused by precipitates.

[0092] Table 2 below shows the size of precipitates measured from surface images of alloys annealed at 750°C and 950°C among the alloys of Preparation Example 1.

[0093] Annealing temperature (°C) Grain boundary precipitates (μm) Intragranular precipitates (μm) Average size (μm) Standard deviation 750 indistinguishable 0.0789 0.385 950 0.5-1 0.02-0.1 0.26 0.113

[0095] Measurement Example 1: Measurement of physical properties of the alloy in Preparation Example 1

[0096] Alloys having the composition of Preparation Example 1 are prepared by annealing at 550°C, 750°C, and 950°C for 1 hour. Additionally, underwater quenching is performed after annealing. The three types of alloys are prepared as specimens for tensile strength testing.

[0097] FIG. 8 is a graph showing the results of a tensile test according to Measurement Example 1 of the present invention.

[0098] Tensile tests were conducted on three types of alloy specimens at temperatures of -196°C and room temperature of 25°C, respectively. The yield strength (YS), ultimate tensile strength (UTS), and elongation (EL) resulting from the tensile tests are shown in Table 3 below.

[0099] Test temperature Specimen annealing temperature YS(MPa) UTS(MPa) EL(%) 25℃ 550℃ 600±40 889±25 19.7±1.9 750℃ 870±14 1045±10 14.4±4.5 950℃ 624±11 932±11 19±0.2 -196℃ 550℃ 1367±65 1775±51 19.4±2.4 750℃ 1198±37 1459±29 30.5±3.4 950℃ 1095±42 1404±17 24.3±2.6

[0100] As shown in FIG. 8 and Table 3 above, the alloys of Preparation Example 1 exhibit high yield strength and maximum tensile strength at -196℃. In addition, they also show high values ​​for elongation, which indicates ductility. As of the filing date, Invar 36, which is used as a low-expansion, high-strength alloy, is known to have a yield strength of approximately 240 MPa and a maximum tensile strength of 430 MPa. In comparison, the alloys of Preparation Example 1 exhibit excellent tensile properties.

[0101] FIG. 9 is a graph showing the magnetic properties of alloy specimens according to Measurement Example 1 of the present invention.

[0102] Referring to FIG. 9, graph (a) illustrates magnetization according to the application of an external magnetic field. The application of the magnetic field and the measurement of the magnetization state are carried out at room temperature, 25°C. When determining the magnetization state of the specimens of Preparation Example 1, they exhibit a general ferromagnetic signal, the coercivity is very small, and considering the increase in magnitude according to the magnetic field, they exhibit soft magnetic characteristics. However, the alloy annealed at 950°C shows a relatively high saturation magnetization value. This is interpreted as being due to a change in the matrix composition as the Mo-rich μ phase precipitates at high temperatures.

[0103] Graph (b) shows the zero field cooling (ZFC) magnetization characteristics according to temperature. At 550°C and 750°C, the ZFC magnetization characteristics change with temperature, but the alloy specimen annealed at 950°C is judged to maintain its magnetized state.

[0104] FIG. 10 is a graph showing the coefficient of thermal expansion of alloy specimens according to Measurement Example 1 of the present invention.

[0105] Referring to Fig. 10, the coefficient of thermal expansion is low due to the magnetism of the alloy below the Curie temperature.

[0106] This phenomenon is because thermal expansion due to lattice vibrations and the lattice contraction effect due to spontaneous volume magnetostriction are offset. Therefore, when the ferromagnetic properties disappear above the Curie temperature, the spontaneous volume magnetostriction effect also disappears, so the coefficient of thermal expansion increases.

[0107] The alloy specimens of the present manufacturing example exhibit a coefficient of thermal expansion of 2.4 ppm / ℃ for the alloy annealed at 550℃, 1.9 ppm / ℃ for the alloy annealed at 750℃, and 1.6 ppm / ℃ for the alloy annealed at 950℃ under low temperature conditions of minus 160℃. Additionally, at room temperature, they exhibit a coefficient of thermal expansion of 4.5 ppm / ℃ to 8.6 ppm / ℃ depending on the annealing temperature.

[0108] Considering that the coefficient of thermal expansion of a typical Invar alloy is 1.2 ppm / ℃, it is confirmed that the alloy of the present example has a coefficient of thermal expansion similar to existing Invar alloys and other Invar-based alloys.

[0110] Preparation Example 2: Fe 45.9 Ni 10.3 Co 33.4 Cr 2.5 Mo 7.9 manufacturing

[0111] Metal raw materials of Fe, Ni, Co, Cr, and Mo with a purity of 99% or higher are prepared. The mixing ratio of these metals is 45.9 : 10.3 : 33.4 : 2.5 : 7.9 in at.% units.

[0112] After placing the raw metal material into a ceramic crucible, it is melted at 1300°C to 1400°C for 4 minutes using induction heating, and then cast using a carbon mold.

[0113] The homogenization heat treatment, cold rolling, annealing, and underwater quenching processes for the casting are the same as those disclosed in Manufacturing Example 1 above.

[0114] Figure 11 is a graph showing the change in hardness according to the annealing temperature of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention.

[0115] Referring to Fig. 11, the sample annealed at 550°C exhibits the highest hardness. Additionally, high hardness is exhibited in the annealing temperature range of 550°C to 650°C, and the hardness of the alloy tends to decrease when the temperature exceeds this range. When the temperature rises in the annealing temperature range above 550°C, recrystallization due to annealing proceeds, and the hardness of the alloy decreases.

[0116] Below, in order to determine the tendency of precipitates within the alloy, the annealing conditions are set to 550°C, 750°C, and 950°C, just as in Preparation Example 1. Subsequently, annealing is performed at different times to check the change in hardness according to the annealing time under the set annealing conditions.

[0117] FIG. 12 is a graph showing the change in hardness according to the annealing time of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention.

[0118] Referring to Fig. 12, the alloy exhibits the highest hardness when the annealing time is 1 hour. During the initial stages of annealing, the hardness temporarily decreases mainly due to recovery, but then increases again as precipitates are formed, reaching the highest hardness at 1 hour. After 1 hour, the hardness gradually decreases as precipitates coarsen and the grains recrystallize.

[0119] XRD analysis results for the alloys of Preparation Example 2 above confirm that they all have a face-centered cubic structure. This is the same as described in Preparation Example 1.

[0120] Figure 13 shows FE-SEM BSE (Field Emission-Scanning Electron Microscope Backscattered Electrons) images according to the annealing temperature of the low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention.

[0121] Referring to FIG. 13, images (a) and (b) show an alloy annealed at 550°C. In the images, a plate-like σ phase with orientation is shown by cold rolling.

[0122] Additionally, images (c) and (d) are images of an alloy annealed at 750°C, and spherical μ phases are observed within a linear σ phase. However, the matrix does not undergo recrystallization, and the precipitate μ phase is observed around the σ phase.

[0123] Images (e) and (f) are images of an alloy annealed at 950°C. In the alloy, recrystallization proceeds and grown grains are observed, and numerous μ phases are precipitated at the grain boundaries. In addition, fine-sized μ phases are also precipitated within the grains.

[0124] Figure 14 is an XRD graph analyzing the crystal structure of precipitates in a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention.

[0125] Referring to FIG. 14, graph (a) is an XRD graph of an alloy that has been annealed at 550°C for 1 hour. In the graph, a peak of the σ phase precipitate is observed. That is, a σ phase with a tetragonal structure is observed.

[0126] Graph (b) is an XRD graph of an alloy annealed at 750°C, and the σ phase and μ phase are observed simultaneously. This result is consistent with images (c) and (d) of Fig. 13.

[0127] Graph (c) is an XRD graph of an alloy annealed at 950°C, and multiple peaks of the μ phase are observed. That is, the σ phase, which was stable under relatively low-temperature annealing conditions, becomes thermodynamically unstable under the relatively high-temperature annealing condition of 950°C, while the μ phase becomes a stable phase. Therefore, as the annealing temperature increases, the σ phase that was previously formed or could have been formed is re-dissolved into the matrix, and only the μ phase precipitates.

[0128] The following Table 4 shows the compositional analysis of the matrix and precipitates of the alloy of Preparation Example 2 via Point-EDS.

[0129] annealing temperature type Fe Co Ni Cr Mo 550℃ matrix 46.1±0.2 35.1±0.2 10.2±0.1 1.4±0.1 7.2±0.1 precipitates 42.3±1.4 32.7±0.8 11.3±0.4 2.1±0.3 11.5±1.6 750℃ matrix 45.9±0.1 35.2±0.2 10.1±0.2 1.4±0.1 7.3±0.3 precipitates 42.5±1.4 33.2±0.8 11.0±1.1 1.7±0.1 11.5±1.3 950℃ matrix 46.5±0.2 35.2±0.2 10.0±0.2 1.4±0.1 7.0±0.1 precipitates 37.0±2.2 31.1±1.3 6.0±0.8 1.4±0.1 24.4±4.5

[0130] In Table 4 above, the unit of composition is at.%, and these are alloys annealed for 1 hour. The metallic element whose content ratio in the precipitates relative to the matrix increased at 550°C and 750°C is Cr. That is, Cr and Mo appear abundant in the precipitates. This is because the precipitates are in the σ phase, and the σ phase is based on intermetallic compounds that are abundant in Cr and Mo.

[0131] In addition, the proportion of Mo in the precipitates increases significantly in the alloy annealed at 950°C. This confirms that Mo contributes significantly to the formation of the μ phase. However, a peculiar aspect is that no change in the Cr content is observed. That is, while an increase in Cr in the precipitates relative to the matrix contributes to the formation of the σ phase, the formation of the σ phase is not observed in the alloy annealed at 950°C because there is no change in the Cr content.

[0132] Figure 15 shows the EBSD map of a low thermal expansion high-strength alloy according to Manufacturing Example 2 of the present invention.

[0133] Image (a) is an Inverse Pole Figure (IPF) map, Image (b) is a Misorientation Angle map, and Image (c) is a Kernel Average Misorientation (KAM) map. Additionally, Image (d) is a Grain Orientation Spread (GOS) map.

[0134] Referring to the images in Fig. 15 above, the alloy annealed at 550°C and 750°C does not undergo recrystallization, so large grains are observed that retain the microstructure created by cold rolling.

[0135] In addition, in the alloy annealed at 950℃, recrystallization proceeds by more than 80%, and fine grains with an average size of 10 μm are observed.

[0136] In addition, the investigation and size of the precipitates of the alloys of Manufacturing Example 2 are summarized in the following Table 5.

[0137] Annealing temperature (°C) Grain boundary precipitates (μm) Intragranular precipitates (μm) Average size (μm) Standard deviation 750 indistinguishable 0.11 0.0514 950 0.5-1 0.02-0.1 0.224 0.115

[0139] Measurement Example 2: Measurement of physical properties of the alloy in Preparation Example 2

[0140] Alloys having the composition of Preparation Example 2 are prepared by annealing at 550°C, 750°C, and 950°C for 1 hour. Additionally, underwater quenching is performed after annealing. The three types of alloys are prepared as specimens for tensile strength testing.

[0141] FIG. 16 is a graph showing the results of a tensile test according to Measurement Example 2 of the present invention.

[0142] Tensile tests were conducted on three types of alloy specimens at temperatures of -196°C and room temperature of 25°C, respectively. The yield strength (YS), ultimate tensile strength (UTS), and elongation (EL) resulting from the tensile tests are shown in Table 6 below.

[0143] Test temperature Specimen annealing temperature YS(MPa) UTS(MPa) EL(%) 25℃ 550℃ 1076±30 1243±29 5.9±1.3 750℃ 792±24 1026±23 18.5±3.6 950℃ 1060±10 1184±10 6.8±3.6 -196℃ 550℃ 1366±41 1376±35 5.6±0.7 750℃ 839±194 1014±187 8.9±4.9 950℃ 1298±21 1421±21 22.6±1.1

[0144] Referring to Table 6 above, it is confirmed that at -196℃, the specimens of Preparation Example 2 showed reduced yield strength and maximum tensile strength compared to the specimens of Preparation Example 1. Since the specimens of Preparation Example 2 had a reduced Mo content compared to Preparation Example 1, it is confirmed that the strength of the alloy improves as the amount of added Mo increases. In addition, the elongation of the specimens of Preparation Example 2 decreased overall compared to the specimens of Preparation Example 1.

[0145] However, the specimens of Preparation Example 2 show a tendency for elongation to increase with increasing annealing temperature in a tensile test at -196℃. This is a phenomenon caused by recrystallization due to the increase in annealing temperature.

[0146] FIG. 17 is a graph showing the magnetic properties of alloy specimens according to Measurement Example 2 of the present invention.

[0147] Referring to FIG. 17, graph (a) illustrates magnetization according to the application of an external magnetic field. The application of the magnetic field and the measurement of the magnetization state are carried out at room temperature, 25°C. The specimens of Preparation Example 2 show relatively higher coercivity compared to the specimens of Preparation Example 1.

[0148] Graph (b) shows the magnetization characteristics according to temperature. All specimens of Preparation Example 2 maintain a magnetized state at room temperature, and are predicted to have a Curie temperature at which magnetization disappears above room temperature. Considering the relationship between the magnetization arrangement of the alloy and the coefficient of thermal expansion, the specimens of Preparation Example 2 are predicted to have a relatively lower coefficient of thermal expansion compared to the specimens of Preparation Example 1.

[0149] FIG. 18 is a graph showing the coefficient of thermal expansion of alloy specimens according to Measurement Example 2 of the present invention.

[0150] Referring to Fig. 18, the coefficient of thermal expansion exhibits low values ​​due to the magnetism of the alloy below the Curie temperature. In particular, under low temperature conditions of minus 160°C, the 550°C annealed specimen shows a low coefficient of thermal expansion of 1.5 ppm / °C, the 750°C annealed specimen shows 1.6 ppm / °C, and the 950°C annealed specimen shows 1.55 ppm / °C. These values ​​are similar to those of Invar alloy, confirming a very low coefficient of thermal expansion.

[0151] In addition, the coefficient of thermal expansion tends to increase due to the removal of magnetization under conditions exceeding the Curie temperature. At room temperature, alloys exhibit a coefficient of thermal expansion of 4.2 ppm / ℃ to 7.0 ppm / ℃, and as the annealing temperature increases, they show a lower coefficient of thermal expansion, which is a phenomenon attributed to differences in the matrix composition caused by precipitates.

[0153] Preparation Example 3: Fe 55.3 Ni 20.1 Co 18.6 Cr 1.4 Mo 4.6 manufacturing

[0154] Metal raw materials of Fe, Ni, Co, Cr, and Mo with a purity of 99% or higher are prepared. The mixing ratio of these metals is 55.3 : 20.1 : 18.6 : 1.4 : 4.6 in at.% units.

[0155] After placing the raw metal material into a ceramic crucible, it is melted at 1300°C to 1400°C for 4 minutes using induction heating, and then cast using a carbon mold.

[0156] The casting is heat-treated at 1200°C for 6 hours to homogenize it. Then, the homogenized alloy is cold-rolled with a reduction rate of 50%.

[0157] The above rolled alloys are annealed at 550°C, 650°C, 750°C, 850°C, 950°C, and 1050°C, respectively. The annealing time is set to 1 hour. Afterward, rapid cooling in water is performed to obtain a low thermal expansion high-strength alloy.

[0159] Preparation Example 4: Fe 62.5 Ni 29.3 Co 6.8 Cr 0.5 Mo 0.9 manufacturing

[0160] Metal raw materials of Fe, Ni, Co, Cr, and Mo with a purity of 99% or higher are prepared. The mixing ratio of these metals is 62.5 : 29.3 : 6.8 : 0.5 : 0.9 in at.% units.

[0161] After placing the raw metal material into a ceramic crucible, it is melted at 1300°C to 1400°C for 4 minutes using induction heating, and then cast using a carbon mold.

[0162] The casting is homogenized at 1200°C for 6 hours. Then, the homogenized alloy is cold-rolled with a reduction rate of 50%.

[0163] The above rolled alloys are annealed at 550°C, 650°C, 750°C, 850°C, 950°C, and 1050°C, respectively. The annealing time is set to 1 hour. Afterward, rapid cooling in water is performed to obtain a low thermal expansion high-strength alloy.

[0165] FIG. 19 is a graph showing the hardness of alloys according to Manufacturing Examples 3 and 4 of the present invention.

[0166] Referring to FIG. 19, the alloys of Manufacturing Examples 3 and 4 exhibit the highest hardness under annealing conditions of 550°C. However, to compare physical properties according to annealing temperature, alloys annealed at 550°C, 750°C, and 950°C are selected for comparison.

[0167] As mentioned above, the hardness of the alloy decreases as the annealing temperature increases. This is a result of the recrystallization that proceeds with the increase in annealing temperature.

[0168] In addition, it is confirmed that the hardness of the alloy in Example 3 is superior to that of Example 4.

[0169] FIG. 20 is a graph comparing the hardness of the commercial Invar 36 alloy of the present invention and the alloys of Manufacturing Examples 3 and 4.

[0170] Referring to FIG. 20, the alloy of Preparation Example 3 exhibits higher hardness than the Invar 36 alloy across all annealing temperatures. However, the alloy of Preparation Example 4 exhibits hardness similar to the existing Invar 36 alloy, or exhibits higher hardness than the Invar 36 alloy when the annealing temperature increases.

[0171] Figure 21 is an XRD graph analyzing the crystal structures of Manufacturing Examples 3 and 4 of the present invention.

[0172] Referring to FIG. 21, graph (a) illustrates the crystal structure of the alloys of Preparation Example 3. Regardless of the annealing temperature, it is confirmed that the matrix of the three alloys of Preparation Example 3 has a face-centered cubic structure.

[0173] Graph (b) illustrates the crystal structure of the alloys of Manufacturing Example 4. The matrix of the alloys has a face-centered cubic structure.

[0174] Figure 22 shows FE-SEM BSE images of the alloys of Manufacturing Examples 3 and 4 of the present invention according to the annealing temperature.

[0175] Referring to FIG. 22, images (a), (b) and (c) show that the alloys of Preparation Example 3 were annealed at 550°C, 750°C, and 950°C, respectively.

[0176] In the alloy annealed at 550°C, it is confirmed that grains with a specific crystal orientation are formed in the rolling direction due to the influence of cold rolling. Additionally, in the alloy annealed at 750°C, grains resulting from partial recrystallization are observed along with linear grains. In the alloy annealed at 950°C, the directional grains caused by cold rolling completely disappear, and entirely new grains appear. In other words, a complete recrystallization phenomenon is observed through annealing at high temperatures.

[0177] In addition, images (d), (e) and (f) show images of the alloys of Preparation Example 4 according to annealing temperatures of 550°C, 750°C, and 950°C.

[0178] Under annealing conditions of 550°C, recrystallization does not proceed, and directional grains resulting from cold rolling are observed. However, in the alloy under annealing conditions of 750°C, recrystallization is confirmed to proceed, and under annealing conditions of 950°C, complete recrystallization proceeds and grain growth is confirmed.

[0180] Measurement Example 3: Measurement of physical properties of alloys in Manufacturing Examples 3 and 4

[0181] FIG. 23 is a graph showing the tensile test results of alloy specimens provided in Manufacturing Examples 3 and 4 according to Measurement Example 3 of the present invention.

[0182] Referring to FIG. 23, the specimens of Manufacturing Example 3 exhibit higher maximum tensile strength and excellent ductility compared to the specimens of Manufacturing Example 4.

[0183] FIG. 24 is a graph showing the coefficients of thermal expansion of alloy specimens of Manufacturing Examples 3 and 4 measured according to Measurement Example 3 of the present invention.

[0184] Referring to FIG. 24, the specimens of Preparation Example 4 have a very low coefficient of thermal expansion at -196°C, which is lower than that of the specimens of Preparation Example 3. This phenomenon can be interpreted as being due to the saturation magnetization value approaching the Invar effect range as a result of the increase in Ni content compared to other preparation examples.

[0185] That is, the specimens of Preparation Example 4 are confirmed to have a coefficient of thermal expansion similar to Invar 36, and actually have a lower value at room temperature.

[0187] Preparation Example 5: Fe 42.91 Ni 10.3 Co 31.0 Cr 3.8 Mo 9.7 V 2.0 C 0.29 manufacturing

[0188] Metal raw materials of Fe, Ni, Co, Cr, Mo, and V with a purity of 99% or higher are prepared. Additionally, raw materials of C are prepared. The sources of the above C are very diverse, and the injection of carbon powder is one method. The mixing ratio of these metals is 42.91 : 10.3 : 31.0 : 3.8 : 9.7 : 2.0 in at.% for Fe, Ni, Co, Cr, Mo, and V, respectively, and carbon is contained at 0.29 at%.

[0189] After placing the raw metal material into a ceramic crucible, it is melted at 1300°C to 1400°C for 4 minutes using induction heating, and then cast using a carbon mold.

[0190] The casting is heat-treated at 1200°C for 6 hours to homogenize it. Then, the homogenized alloy is cold-rolled with a reduction rate of 50%.

[0191] The above rolled alloys are annealed at 550°C, 650°C, 750°C, 850°C, 950°C, and 1050°C, respectively. The annealing time is set to 1 hour. Afterward, rapid cooling in water is performed to obtain a low thermal expansion high-strength alloy.

[0192] The alloy of Manufacturing Example 5 above is prepared by partially reducing the atomic percentage of Co in the composition of the alloy of Manufacturing Example 1 and adding V and C in the reduced amount, so it needs to be compared with the alloy of Manufacturing Example 1 to confirm the effect of adding V and C.

[0193] FIG. 25 is a graph showing the hardness according to the annealing temperature of alloys according to Manufacturing Example 1 and Manufacturing Example 5 of the present invention.

[0194] Referring to FIG. 25, the alloys of Preparation Example 1 and Preparation Example 5 exhibit the highest hardness at an annealing temperature of 550°C. Subsequently, as the annealing temperature increases, a decrease in hardness is observed.

[0195] In addition, it is confirmed that the alloys of Preparation Example 5 have higher hardness compared to Preparation Example 1. This indicates that the hardness of the alloy has increased due to the addition of V and C.

[0196] Below, to determine the temperature-dependent trends of the alloys, measurements are performed by setting the annealing temperature to 550℃, 750℃, and 950℃.

[0197] FIG. 26 is a graph showing the hardness according to the annealing time of alloys according to Manufacturing Example 1 and Manufacturing Example 5 of the present invention.

[0198] Referring to FIG. 26, alloys annealed at 550°C exhibit similar hardness under annealing conditions of 1 hour and 2 hours. At other annealing temperatures, the highest hardness is exhibited under annealing conditions of 1 hour. Below, annealing conditions of 1 hour are applied for the comparison of physical properties.

[0199] FIG. 27 is an XRD graph analyzing the crystal structure of alloys according to Manufacturing Example 5 of the present invention.

[0200] Referring to Fig. 27, the matrix is ​​confirmed to have a face-centered cubic structure under all three annealing conditions. However, precipitates are observed in the alloy annealed at 950°C.

[0201] FIG. 28 is FE-SEM BSE images of alloys according to Manufacturing Example 5 of the present invention.

[0202] Referring to FIG. 28, images (a) through (c) show an alloy that has not been annealed. Images (d) through (f) show an alloy annealed at 550°C, images (g) through (i) show an alloy annealed at 750°C, and images (j) through (l) show an alloy annealed at 950°C.

[0203] In the alloy of Preparation Example 5 with added V, a large amount of precipitates is observed even before annealing is performed. In addition, the size of the precipitates is 3 μm to 5 μm, and the size of the carbides increases compared to other preparation examples.

[0204] In particular, under annealing conditions of 750°C and 950°C, the formation of secondary precipitates is observed in addition to primary coarse precipitates. The secondary precipitates have an average size of 1 μm or less.

[0205] FIG. 29 is an EDX (Energy Dispersive X-ray Spectroscopy) graph analyzing the components of the precipitate according to Preparation Example 5 of the present invention.

[0206] Referring to FIG. 29, graph (a) shows the alloy without homogenization annealing. The remaining graphs (b) to (d) show the composition of precipitates under annealing conditions of 550°C, 750°C, and 950°C. Considering that the quantitative accuracy of carbon (C) in EDX quantitative analysis is low, the results of this analysis include only elements excluding carbon (Fe, Co, Ni, Cr, Mo, etc.).

[0207] An increase in Mo content is observed in the precipitates compared to the components constituting the matrix. In other words, it is confirmed that supersaturated Mo within the matrix diffused and nucleated to form precipitates.

[0208] FIG. 30 is an EBSD IPF (Inverse Pole Figure) map and an Inverse Pole Figures photograph in the RD direction of alloys according to Manufacturing Example 5 of the present invention.

[0209] Referring to Fig. 30, the inverse pole figure (IPF) indicates which direction of the crystal lattice within the alloy the rolling direction (RD) corresponds to.

[0210] Images (a) and (d) show an alloy annealed at 550°C, images (b) and (e) show an alloy annealed at 750°C, and images (c) and (f) show an alloy annealed at 950°C.

[0211] In alloys annealed at 550°C and 750°C, recrystallization does not proceed, so the orientation of the grains in a specific direction is observed. On the other hand, in alloys annealed at 950°C, the grain size increased due to recrystallization, and the orientation of the crystals shows a random tendency.

[0213] Measurement Example 4: Measurement of physical properties of the alloys of Preparation Example 5

[0214] Alloys having the composition of Preparation Example 5 are prepared by annealing at 550°C, 750°C, and 950°C for 1 hour. Additionally, underwater quenching is performed after annealing. The three types of alloys are prepared as specimens for tensile strength testing. Separately, alloys having the composition of Preparation Example 1 are also prepared as specimens for comparison.

[0215] FIG. 31 is a graph showing the results of a tensile test at room temperature according to Measurement Example 4 of the present invention.

[0216] Referring to FIG. 31, the solid lines represent the specimens of Preparation Example 5, and the dotted lines represent the specimens of Preparation Example 1. In either case, it is confirmed that the yield strength and maximum tensile strength of the specimens of Preparation Example 5 have increased significantly compared to Preparation Example 1. This is confirmed by the following Table 7, which quantifies the yield strength and maximum tensile strength of the specimens of Preparation Example 5 in FIG. 32.

[0217] 550℃ 750℃ 950℃ YS(MPa) 1416.9±29.6 1002.8±20.1 635.5±2.6 UTS(MPa) 1505.6±28.3 1191.8±27.2 976.2±0.9

[0218] FIG. 32 is a graph showing the results of a tensile test at -196℃ according to Measurement Example 4 of the present invention.

[0219] Referring to Fig. 32, the yield strength and maximum tensile strength increase significantly due to the addition of V and C. The above results are summarized in the following Table 8.

[0220] 550℃ 750℃ 950℃ YS(MPa) 1730.2±52.0 1386.7 980.4 UTS(MPa) 1772.9±65.3 1569.3 1408.5

[0222] Thermodynamic simulation example

[0223] Thermodynamic simulations are performed on the alloys of Preparation Example 1 and Preparation Example 2. Through thermodynamic simulations, the types of thermodynamically stable phases and the fractions of each phase under specific temperature conditions of Preparation Example 1 and Preparation Example 2 can be calculated. In addition, the composition and crystal structure constituting the matrix and precipitates can be predicted.

[0224] FIG. 33 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 1 according to a simulation example of the present invention.

[0225] Referring to FIG. 33, the melting point of the alloy is formed at less than 1400°C. An austenite structure is formed starting from 500°C, and above 800°C, the austenite structure forms the entire matrix. That is, the matrix is ​​predicted to have a face-centered cubic structure, which is exactly consistent with the XRD graph of Preparation Example 1.

[0226] In addition, the ferrite structure having a body-centered cubic structure is predicted to disappear completely at about 750°C, and the σ phase is also predicted to disappear completely at 800°C or higher.

[0227] FIG. 34 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 1 according to a simulation example of the present invention.

[0228] Referring to Fig. 34, it is observed that changes in the composition of metal elements within the σ phase occur with increasing annealing temperature. As the annealing temperature increases, the Mo content within the σ phase tends to increase, while the Cr and Ni content tends to decrease.

[0229] However, when the annealing temperature exceeds 750°C, the compositional ratio of metal elements is predicted to have a constant trend. That is, in the annealing range up to 800°C, the change in the composition of the σ phase is not significant, and it is predicted to maintain a somewhat stable form.

[0230] FIG. 35 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 2 according to a simulation example of the present invention.

[0231] Referring to Fig. 35 above, since the alloy of Manufacturing Example 2 does not show a significant change in composition compared to Manufacturing Example 1, a change in crystal structure similar to that of Fig. 34 is predicted.

[0232] FIG. 36 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 2 according to a simulation example of the present invention.

[0233] Referring to Fig. 36, the degree of change in the constituent metal within the σ phase shows a trend similar to that of Fig. 34. However, the Cr content is predicted to increase compared to Fig. 34, which suggests that the formation of the σ phase may increase compared to Fig. 34.

[0234] FIG. 37 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 3 according to the simulation example of the present invention.

[0235] Referring to Fig. 37 above, since the alloy of Manufacturing Example 3 does not show a significant change in composition compared to Manufacturing Example 1, a change in crystal structure similar to that of Fig. 34 is predicted.

[0236] However, the formation of the body-centered cubic ferrite phase is weak and is predicted to disappear rapidly depending on the annealing temperature.

[0237] In addition, Figure 37 predicts that the Laves phase will be formed at temperatures below 700°C.

[0238] FIG. 38 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 3 according to a simulation example of the present invention.

[0239] Referring to Fig. 38, the composition of metal elements in the Laves phase shows a different pattern from the metal composition in the σ phase in other manufacturing examples. In addition, no significant change in composition is observed at the annealing temperature.

[0240] FIG. 39 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 4 according to the simulation example of the present invention.

[0241] Referring to Fig. 39, only austenite and ferrite are predicted within the matrix, and the body-centered cubic ferrite is predicted to disappear below 600°C.

[0242] In addition, due to the low concentration of Mo, no precipitates are formed, and it is predicted that only a matrix with low thermal expansion characteristics will be generated.

[0243] FIG. 40 is a graph predicting the crystal structure of the matrix according to the composition of the alloy of Manufacturing Example 5 according to the simulation example of the present invention.

[0244] Referring to Fig. 40, the formation of σ phase and ferrite within the matrix is ​​predicted. At an annealing temperature of 750°C or higher, ferrite is predicted to disappear from the matrix, but the fraction of the σ phase within the matrix gradually decreases as the temperature increases.

[0245] FIG. 41 is a graph predicting the formation of precipitates according to the composition of the alloy of Manufacturing Example 5 according to a simulation example of the present invention.

[0246] Referring to Fig. 41, when the annealing temperature in the σ phase is raised to about 750°C, the amounts of Co and Ni decrease slightly, while the amount of V remains constant or increases slightly. Also, referring to Fig. 40, it is predicted that the amount of precipitate will increase.

[0248] Composition Design Example: Design of Carbide Precipitate Composition and Prediction of Precipitates in Alloys

[0249] In this design example, C is added to the FeNiCoCrMo alloy of the conventional manufacturing example, and C is added to the FeNiCoCrMoV alloy. The formation of carbide is predicted with the addition of C.

[0250] FIG. 42 shows Fe according to a composition design example of the present invention. 42.91 Ni 10.3 Co 31 Cr 3.8 Mo 9.7 C 0.29 This is a graph predicting the phase change of the alloy.

[0251] Referring to Fig. 42, the formation of carbide particles is predicted compared to other manufacturing examples due to the influx of C.

[0252] The influx of C leads to the formation of the carbide phase M6C, which contributes to the improvement of the alloy's strength. That is, at an annealing temperature of 750°C or higher, the ferrite disappears, and a small amount of σ phase and carbide may appear as precipitates.

[0253] FIG. 43 shows Fe according to a composition design example of the present invention 39.81 Ni 9.73 Co 32.93 Cr2.33 Mo 13.00 V 2.0 C 0.2 This is a graph predicting the phase change of the alloy.

[0254] Referring to Fig. 43, a high fraction of the σ phase is observed in the existing V-added alloy due to the addition of V and C, and the formation of the M6C phase of the carbide is predicted.

[0255] According to the present invention described above, strength enhancement is maximized through a multi-component alloy, and the coefficient of thermal expansion exhibits characteristics similar to those of conventional Invar alloys. To enhance strength, fine nano-sized precipitates are formed within a face-centered cubic matrix. These precipitates are formed at grain boundaries or within grains through homogenization heat treatment and annealing processes; while their size or shape may vary depending on the heat treatment conditions, they possess a fine size that is not coarsenic. Depending on the heat treatment conditions, partial recrystallization occurs, creating a heterostructure microstructure. That is, a combination of the heterostructure microstructure and nano-precipitates ensures high strength and excellent ductility.

[0256] Furthermore, the formed face-centered cubic matrix exhibits low thermal expansion characteristics through the multi-component compositional design of the constituent elements, driven by thermal expansion due to lattice vibrations and lattice contraction effects caused by spontaneous volume magnetostriction; additionally, the fine precipitates generated internally do not significantly reduce the coefficient of thermal expansion of the alloy. Consequently, a coefficient of thermal expansion similar to that of conventional Invar alloys can be obtained. Explanation of the symbols

[0257] doesn't exist

Claims

Claim 1 A low thermal expansion high-strength alloy characterized by comprising Fe: 43.2 at.% to 55.3 at.%, Co: 18.6 at.% to 33.4 at.%, Ni: 10.3 at.% to 20.1 at.%, Cr: 1.4 at.% to 3.8 at.%, Mo: 4.6 at.% to 9.7 at.%, and other unavoidable impurities, wherein the mixing ratio of Ni is greater than the mixing ratio of Cr, having a face-centered cubic matrix structure through annealing, containing σ phase or μ phase precipitates within the grain boundaries, and having a coefficient of thermal expansion of 4.2 ppm / ℃ to 8.6 ppm / ℃ at room temperature. Claim 2 A low thermal expansion, high strength alloy according to claim 1, characterized in that the body-centered cubic structure is converted to the face-centered cubic structure through the annealing. Claim 3 A low thermal expansion, high-strength alloy according to claim 1, characterized in that the σ phase has an increased Cr content compared to the base. Claim 4 A low thermal expansion, high strength alloy according to claim 1, characterized in that the μ phase has an increased Mo content compared to the base. Claim 5 A low thermal expansion high-strength alloy according to claim 1, characterized in that the annealing is performed at a temperature of 550°C to 750°C. Claim 6 A low thermal expansion, high-strength alloy according to claim 1, characterized in that the coefficient of thermal expansion decreases as the Mo content decreases. Claim 7 A low thermal expansion high-strength alloy according to claim 1, characterized in that the low thermal expansion high-strength alloy contains C in an amount of 0.29 at.% or less and V is added in an amount of 0 at.% to 2.0 at.%. Claim 8 A low thermal expansion, high strength alloy according to claim 7, characterized in that M6C-phase carbides are included in the precipitates by the addition of C. Claim 9 A low thermal expansion high strength alloy according to claim 1, characterized in that the low thermal expansion high strength alloy further comprises V, and V is added in an amount of 0 at.% to 2.0 at.%. Claim 10 A low thermal expansion, high strength alloy according to claim 9, characterized in that the amount of the precipitate is further increased by the addition of V. Claim 11 In claim 9, the low thermal expansion high strength alloy is characterized by having primary precipitates coarsen through annealing and secondary precipitates smaller than the primary precipitates. Claim 12 A low thermal expansion, high-strength alloy according to claim 11, characterized in that the secondary precipitates have an average size of 1 μm or less. Claim 13 A low thermal expansion high strength alloy according to claim 9, characterized in that the low thermal expansion high strength alloy further comprises C, and the C is added in an amount of 0 at.% to 0.29 at.%. Claim 14 A low thermal expansion, high strength alloy according to claim 13, characterized in that M6C-phase carbides are included in the precipitates by the addition of C. Claim 15 A step of performing a homogenization heat treatment on a casting formed by melting and casting a metal raw material; a step of performing cold rolling after the homogenization heat treatment; A method for manufacturing a low thermal expansion high-strength alloy, comprising the step of forming a low thermal expansion high-strength alloy by controlling the grains and precipitates forming the matrix of the alloy through annealing of the cold-rolled alloy, wherein the low thermal expansion high-strength alloy comprises Fe: 43.2 at.% to 55.3 at.%, Co: 18.6 at.% to 33.4 at.%, Ni: 10.3 at.% to 20.1 at.%, Cr: 1.4 at.% to 3.8 at.%, Mo: 4.6 at.% to 9.7 at.%, and other unavoidable impurities, wherein the mixing ratio of Ni is greater than the mixing ratio of Cr, and through annealing, has a matrix having a face-centered cubic structure, contains σ-phase precipitates or μ-phase precipitates within the grain boundaries, and has a coefficient of thermal expansion of 4.2 ppm / ℃ to 8.6 ppm / ℃ at room temperature. Claim 16 A method for manufacturing a low thermal expansion high-strength alloy, characterized in that, in claim 15, the annealing is performed at 550°C to 950°C. Claim 17 A method for manufacturing a low thermal expansion high-strength alloy according to claim 16, characterized in that the size of the crystal grains increases through recrystallization of the crystal grains at an annealing temperature of 950℃. Claim 18 delete Claim 19 A method for manufacturing a low thermal expansion high-strength alloy according to claim 15, characterized in that the above-mentioned precipitate is generated by heat treatment and has a chemical composition or crystal structure different from the above-mentioned base. Claim 20 A method for manufacturing a low thermal expansion high-strength alloy, wherein, in claim 15, when the annealing temperature rises, a TCP (Topologically Close-Packed) precipitate is formed within the grain boundaries.

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