KR20130142467A - Titanium-based bulk amorphous matrix composite and method of fabricating thereof - Google Patents

Abstract

An in-situ Ti-based bulk amorphous matrix composite has a composition (at%) represented by the below chemical formula 1. [The chemical formula 1] Ti_aZr_bBe_cCu_dNi_eM_fI_g In the chemical formula 1, M is one or more of Nb or Ta, and I is an impurity, and 38<=a<=50, 11<=b<=18, 12<=c<=20, 6<=d<=10, 6<=e<=9, 1<=f<=20, and 0.01<=g<=0.5. Moreover, a + b + c + d + e + f + g = 100.

Description

Titanium-based bulk amorphous matrix composite and method for manufacturing the same {TITANIUM-BASED BULK AMORPHOUS MATRIX COMPOSITE AND METHOD OF FABRICATING THEREOF}

Titanium based amorphous matrix composites and methods for their preparation are provided.

Due to energy depletion and global warming, alternative energy development is under way. Since hydrogen is abundant, there is a growing interest in hydrogen energy that can be used in fuel cells among sustainable and clean alternative energy. Hydrogen must be extracted from other gases or compounds before it can be used as fuel. Although there are several methods such as chemical or electrolytic methods for hydrogen extraction, permeation using a solid metal membrane has high selectivity, high yield, and a compact and simple process. It is widely used to produce high purity hydrogen.

The principle of hydrogen permeation is to heat a mixed gas such as "CO + H ₂ " to decompose hydrogen gas molecules into hydrogen protons on the surface of the membrane, diffuse hydrogen protons through the membrane's atomic structure, and Recombination into molecules. "CO + H ₂ " can be obtained from steam reforming of methane gas at about 400 degrees Celsius. Hydrogen permeation through a solid membrane follows Fick's first law and is as follows.

Where J is the flux, which means the amount of hydrogen moving per unit time per unit area, D is the hydrogen diffusion coefficient, and δC / δx is between the permeate side and the retentate side of the membrane. Concentration gradient.

Fick's first law can also be expressed as

Is the difference in hydrogen concentration between the permeate side and the residue side relative to the thickness t of the membrane. The larger the hydrogen flux value is the higher the yield of pure hydrogen, which can be achieved with a thinner film. However, the thickness of the film is preferably thick enough to withstand the hydrogen pressure applied at the inlet side.

Where K _s is the material constant and P _H2 is the hydrogen partial pressure.

Hydrogen flux is expressed as follows.

The hydrogen flux varies with the pressure difference between the permeate side and the residue side of the membrane, the permeability and solubility of hydrogen, and the thickness of the membrane.

Palladium and its alloys, particularly palladium-silver alloys, are the material of solid membranes, and have excellent hydrogen permeability, durability, resistance to H ₂ O, CO, CO ₂ and H ₂ S, and hydrogen membranes in the United States. (hydrogen membrane), gas separation apparatus (gas separation apparatus) and the like. However, due to the scarcity and high price of palladium, alternative materials that can be used in hydrogen separation membranes instead of palladium are being actively developed.

As a substitute for palladium, various kinds of alloys having a crystalline or amorphous structure have been proposed. The first amorphous glass was made nearly 50 years ago using the melt spinning technique by Caltech's team. Unlike nonmetal amorphous solids such as ceramic glass, metal alloys have excellent mechanical properties in the amorphous state. The amorphous state is excellent in physical, chemical, and mechanical properties because atoms are aperiodic in arrangement, and thus can be applied to various fields. With regard to elasticity, the stiffness of the amorphous metal alloy is about 30% lower than that of the crystalline metal alloy, and the yield strength of the amorphous metal alloy is about 2-3 times greater than that of the crystalline metal alloy. However, the amorphous metal may be deformed when the elastic strain is large. In the 1970s Fe-P (6.0 wt%)-C (1.7 wt%) amorphous metals were commercialized for the first time due to their soft magnetism and were used to improve the efficiency of transformers and magnetic heads. It became. The amorphous metal of Fe-P-C was developed in 1967 by researchers at Caltech. In addition, the vacuum casting technique has made it possible to produce amorphous metals in bulk, for example at low cooling rates. New types of alloys have been developed, and nowadays bulk materials with diameters of at least about 1 cm contain rare earth elements and elements such as Mg, Ca, Zr, Pd, Ti, Cu, Fe, Co, Ni It is made from the field. However, Al may be made into a rod shape having a diameter of about 1 mm or less. Although excellent physical, chemical and mechanical properties can be achieved by these various types of alloys, the application of amorphous metals has been limited for several reasons. First, amorphous metals lack the reliability and reproducibility of physical properties. Because the alloys thus far developed are sensitive to impurities, especially oxygen, which causes severe embrittlement. Second, these alloys have inherently fragile properties associated with oxide glass. For example, such alloys cannot be plastically deformed at ambient temperature in tensile mode.

Research on amorphous metals is rapidly advancing, and the next generation of amorphous metals may exhibit ductile behavior and may be applied in new fields. Applications include corrosion-resistant coatings, MEMS, biomedical applications, and other applications such as ships, auto parts, bicycles, tennis rackets and golf clubs.

However, compared with crystalline alloys, the development level of amorphous metals is still very small, and the exact mechanism for controlling the properties such as plastic deformation is still unknown. At low temperatures, when stress is greater than yield stress, plastic deformation occurs by localized atomic flow in the thin shear band. Recent investigations have shown that shear bands in amorphous metals are caused by structural softening rather than by thermo-mechanical softening of the material. Structure seems to play an important role in the deformation mechanism. Despite the acceptance of the free-volume model proposed by Spaepen, a local shear transformation zone (STZ) mechanism that emphasizes atomic flow in viscous material is obtained by molecular dynamic simulation. It tends to be supported by the results.

Inelastic deformation in amorphous metals has been interpreted as local shearing of approximately 30-50 element clusters with strain-softening and further inelastic dilatation to form shear bands. . As the shear strain increases, the fully strained STZ affects the surrounding material and helps to produce thin shear bands approximately 10-15 nm thick.

The plastic deformation mechanism in amorphous metals is fundamentally different from the plastic deformation mechanism in crystalline solids. This is because the atomic structure of amorphous metals has very little long-range order. In tension mode the shear bands do not stretch and remain local and unstable. When the plastic strain value is very large in the thin region, point defects accumulate that cause crack formation and brittle fracture.

Increasing the plastic deformation of amorphous metals in low temperature tensile mode introduces inhomogeneities into the amorphous structure, which can generate a large number of STZs and hinder the growth of shear bands. Although some form of heterogeneity can be introduced into the amorphous structure, the dynamic growth of shear bands requires large and close obstacles relative to the thickness of the shear bands. Methods of achieving plastic deformation in tensile mode include ex-situ and in-situ production methods of bulk metallic glass matrix composites. The ex situ method allows for tensile ductility at the expense of strength, but the manufacturing process is complex. The in situ method developed by the Caltech researchers brought the plastic deformation of the Zr-Ti-Nb-Cu-Be in situ BMG matrix composite up to 9% in tension mode. Mechanical behavior is determined by strain softening, and the maximum strength of the ductile in situ BMG matrix composite is reduced by about 15% compared to monolithic zirconium-based BMG. The structure of this complex is determined by the presence of the Nb-rich dendritic phase in the zirconium-based amorphous structure.

This type of in situ BMG matrix composite is similar to the crystalline Nb- (Ti, Zr) -Ni two-phase alloy produced by the Aoki group of Kitami Institue, Japan, and is a crystalline Nb- ( The Ti, Zr) -Ni binary phase alloy is composed of Nb-rich dendrite containing Nb in the Ni- (Ti, Zr) crystalline matrix. Attention has been focused on the emergence of the Nb-rich phase, which indicates that the permeability of hydrogen in the Group 5 element of the periodic table with body centered cubic (bcc) atomic structure is face centered cubic, It is experimentally proved that the permeability of hydrogen in palladium having a fcc) structure is much larger than that of hydrogen. However, for these bcc metals such as V, Nb, Ta and the like, insufficient permeability is not a problem but a low mechanical strength under hydrogen gas such as hydrogen embrittlement.

Amorphous metals have already been used as materials for hydrogen films. Ni-Nb based amorphous alloys or Ni-Zr based amorphous alloys have been demonstrated to exhibit excellent performance with hydrogen permeability values of about 2x10 ^-8 mol.msPa ^1/2 (400 degrees Celsius), similar to the hydrogen permeability values of palladium.

In order for the amorphous alloy to have a much higher hydrogen permeability value than the palladium alloy, a composite structure in which a group 5 element-rich phase is dispersed in the amorphous structure is prepared. For example, the dispersion of the continuous crystalline phase will help rapid diffusion of the hydrogen atoms, and the amorphous matrix will serve to control the lattice expansion of the dispersed phase with hydrogen and to prevent hydrogen embrittlement. .

One embodiment according to the invention is to improve hydrogen permeability.

One embodiment according to the invention is to improve the strength at low temperatures.

And can be used to achieve other tasks not specifically mentioned other than the above tasks.

One embodiment of the present invention is to in-situ crystalline-reinforced Ti-based bulk amorphous matrix composites (BAMC) as a solid membrane for gas separation, such as polymer electrolyte membrane of a fuel cell It's about application. The microstructure of the film formed by vacuum casting technology has a bcc crystalline structure with dendritic morphology located in the Ti-based amorphous matrix.

Composites formed at various composition ratios are relatively easy to make because of their high glass forming ablitiy. For example, the volume ratio and size of reinforcement, ie, the bcc crystalline phase, may be 1-40% and 0.01-100 μm, respectively. The dendritic morphology for this reinforcement can be isolated, thin and can have directionality depending on the required properties.

The composite has excellent properties such as high hydrogen permeability, high mechanical strength, high corrosion resistance and the like. For example, a bcc phase with many Group 5 elements, such as Nb and / or Ta, has a high hydrogen permeability, which is 6x10 ^-8 mol / msPa ^1/2 at 325 degrees Celsius to 375 degrees Celsius, which is the same experiment. 2 to 3 times larger than palladium based alloys under conditions. Hydrogen permeability is an important factor for solid membrane materials used to separate or purify hydrogen. The composite shows high strength and high malleability at room temperature. The composite shows excellent corrosion resistance in 1M sulfuric acid at 80 degrees Celsius with hydrogen gas or air bubbling. The appearance of a second phase improves the mechanical strength in the supercooled liquid region. Because of these relatively inexpensive costs as well as excellent physical, mechanical and chemical properties, in situ BMG matrix composites can be used for the separation and purification of hydrogen.

The Ti-based bulk amorphous matrix composite includes a composition represented by Formula 1 below.

[Formula 1]

Ti _a Zr _b Be _c Cu _d Ni _e M _f I _g

In Formula 1, M is at least one of Nb or Ta, I is impurity, 38≤a≤50, 11≤b≤18, 12≤c≤20, 6≤d≤10, 6≤e≤9, 1 ≤ f ≤ 20, and 0.01 ≤ g ≤ 0.5, and a + b + c + d + e + f + g = 100.

One embodiment according to the present invention can improve hydrogen permeability and improve strength at low temperatures.

1A and 1B are binary phase diagrams of Ti-Nb and Ti-Ta, respectively.
1C, 1D, and 1E are schematic diagrams showing shell structures of Ti, Nb, and Ta, respectively.
1F is a graph showing the principle of in situ complex formation.
2A is a schematic diagram showing a vacuum suction casting mechanism.
2B is a schematic diagram showing a vacuum squeeze casting mechanism.
3A is an XRD graph of Ø10 mm (Ti-Zr-Be-Cu-Ni) ₉₅ Nb ₅ and Ø7 mm (Ti-Zr-Be-Cu-Ni) ₉₅ Ta ₅ bars.
3B is a DSC graph of Ø10 mm (Ti-Zr-Be-Cu-Ni) ₉₅ Nb ₅ and Ø8 mm (Ti-Zr-Be-Cu-Ni) ₉₅ Ta ₅ bars.
3C is an XRD graph of Ø7 and Ø10 mm diameter bars of (Ti-Zr-Be-Cu-Ni) ₉₅ Ta ₅ .
3D and 3E are SEM images of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ (Ø8 mm) and (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ (Ø8 mm), respectively. .
4A is a phase diagram showing the range of formation of Ti-Zr-Be-Cu-Ni-Nb in situ amorphous metal composites.
4B and 4C are SEM photographs showing the microstructures of alloys A2 and A3 containing 10 at% and 15 at% of Nb, respectively.
4D is an XRD graph of alloys A0, A1, A2, A3 containing 0, 5, 10, and 15 at% of Nb, respectively.
4E is a DSC graph of alloys (A1, A2, A3) containing 0, 5, 10, and 15 at% of Nb, respectively.
4F and 4G are SEM images showing the microstructure of Ti ₄₅ (Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) _45/55 Nb ₁₀ and Ti _41.6 Zr _12.6 Be _15.8 Cu _7.9 Ni _7.1 Nb ₁₅ alloy, respectively.
5a is an XRD graph of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) _{100- x} Nb _x alloy (x = 0, 5, 10, 15 at%).
5b is a DSC graph of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) _{100- x} Nb _x alloy (x = 0, 5, 10, 15 at%).
5C, 5D, 5E, and 5F are SEM images of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) _{100- x} Nb _x alloys (x = 0, 5, 10, 15 at%), respectively.
6A is an XRD graph of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ alloy.
6B is a SEM photograph of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ alloy.
6C is a DSC graph of (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ alloy.
FIG. 7A is a stress-strain curve when compressed at a strain rate of 10 ⁻⁴ s ^{− 1} for an in situ Ti-based BMS composite including Nb and Ta. FIG.
Figure 7b is 10 ^-4 s with respect to the in-situ Ti-based BMS complex comprising ^Nb-strain curve is - stress at the time of compression by the strain rate of ^1.
7C, 7D and 7E are stress-strain curves at strain rates of 10 ⁻⁵ s ⁻¹ , 10 ⁻⁴ s ⁻¹ , and 10 ⁻³ s ⁻¹ , respectively.
7F is a graph showing the change in flow stress according to the applied strain rate.
FIG. 8A is a graph showing the change in flux at 350 degrees Celsius for an in situ Ti-based BMG composite comprising 5 and 15 at% Nb prepared by vacuum squeeze casting.
FIG. 8B is a graph showing changes in flux at 350 degrees Celsius and 375 degrees for an in situ Ti-based BMG composite comprising 10 at% Nb prepared by vacuum suction casting.
8C and 8D are graphs showing hydrogen permeability of Ti-BMG composites prepared by vacuum suction casting or vacuum squeeze casting, respectively, compared with Pd-Cu alloy and Ni-based amorphous.
FIG. 8E is a graph showing hydrogen permeability of Ti / Ta-BMG composites prepared by vacuum squeeze casting compared to Ti / Nb BMG composites according to pressure change, and FIG. 8F is Ti / Ta- prepared by vacuum squeeze casting. It is a graph showing the hydrogen permeability of the BMG composite compared with the Pd-Cu alloy and Ni-based amorphous metals according to the change of the inverse of the temperature.
FIG. 9A shows a potentio-dynamic graph of an in situ Ti-BMG composite at 1 M H ₂ SO ₄ and 2 ppm F ⁻ with H ₂ bubbling at 80 degrees Celsius with stainless steel and Ti-6Al-4V. It is a figure compared and shown.
9B shows the corrosion current density and protective current according to the Nb content (%) for in situ Ti-BMG composites at 1 M H ₂ SO ₄ and 2 ppm F ⁻ with H ₂ bubbling at 80 degrees Celsius. It is a graph showing a change in density (passivation current density).
9C is a graph showing the change in corrosion current density over time for a potentio-static test performed under a potential of 1M H ₂ SO ₄ and 2 ppm F ⁻ with H ₂ bubbling at 80 degrees Celsius. to be.

BRIEF DESCRIPTION OF THE DRAWINGS The above and other features and advantages of the present invention will be more apparent from the following detailed description taken in conjunction with the accompanying drawings, in which: FIG. The present invention may be embodied in many different forms and is not limited to the embodiments described herein. In order to clearly illustrate the present invention, parts not related to the description are omitted, and the same reference numerals are used for the same or similar components throughout the specification. In the case of publicly known technologies, a detailed description thereof will be omitted.

In the drawings, the thickness is enlarged to clearly represent the layers and regions. It will be understood that when an element such as a layer, film, region, plate, or the like is referred to as being "on" another portion, it includes not only the element directly over another element, On the other hand, when a part is "just above" another part, there is no other part in the middle. Conversely, when a part of a layer, film, region, plate, etc. is "below" another part, this includes not only the other part "below" but also another part in the middle. On the other hand, when a part is "just below" another part, it means that there is no other part in the middle.

Next, an embodiment according to the present invention will be described in detail with reference to the accompanying drawings.

According to one embodiment of the present invention, an in-situ bulk metallic glass composite is used as the material of a permeable solid membrane used in gas separation techniques for the separation and purification of hydrogen. Alloys produced by vacuum arc melting technology in combination with vacuum suction casting or vacuum squeeze casting technology exhibit superior properties than pure palladium metals or palladium alloys. Before describing the hydrogen permeability, mechanical and chemical properties, the preparation of a membrane having a microstructure exhibited by a secondary electron microscope (SEM) and X-ray diffraction and thermal calorimetry (DSC) is first described.

In order to find the most suitable industrial applications for the formation of crystalline phases with dendritic structures in amorphous matrices and the physical, chemical and mechanical properties of materials formed in in situ BMG matrix composites.

In the Ti-based amorphous metal matrix thermal mixing of the negative (negative heat of mixing) (DELTA _{H Ti - Be = -30 kJ /} mol, DELTA _{H Ti - Ni = -35 kJ /} mol, DELTA H _{Ti - Cu} = -9 kJ / mol, DELTA _{H Zr - Be = -43 kJ /} mol, DELTA _{H Ti - Cu = -23 kJ /} mol, DELTA _{H Ti - Ni = -49 kJ /} mol) and the atomic size ratio (atom ratio size) ( The matrix is formed of suitable elements with r _Ti / r _Zr = 0.92, r _Cu / r _Ti = 0.87, r _Be / r _Cu = 0.875, r _Be / r _Ni = 0.9). The alloys thus prepared may correspond to compositions close to eutectic compositions with low melting temperatures, wherein the melting temperatures are from about 1000 degrees Celsius to about 1050 degrees Celsius, and are Ti, Zr, Cu, Ni, Be crystalline. The alloys are lower than about 1660 degrees Celsius, about 1852 degrees Celsius, about 1084 degrees Celsius, about 1453 degrees Celsius, and about 1278 degrees Celsius, respectively. Although the amorphous phase may be formed with a relatively low free energy and a high viscosity of the liquid, the amorphous phase is meta stable, which can prevent the diffusion of a wide range of elements so that the cooling rate is critical When larger, crystallization can be prevented.

In order to produce an amorphous metal matrix matrix used as a material for the hydrogen film, a crystalline phase rich in Group 5 elements is formed. For example, as described in Example 2, the formation of the beta phase can be promoted by adding Nb and / or Ta elements to the Ti-based amorphous metal.

In addition to excellent hydrogen permeability properties, in-situ amorphous matrix composites have a high strength of approximately 1.6 to 2 GPa and a low density of approximately 5.1 to 5.5 g / cm ³ , high formability, and high corrosion resistance.

According to one embodiment of the invention, the Ti-based in-situ bulk amorphous matrix composite has a composition (at%) represented by the following formula (1).

[Formula 1]

Ti _a Zr _b Be _c Cu _d Ni _e M _f I _g

In Formula 1, M is at least one of Nb or Ta, I is an impurity such as oxygen or carbon, 38≤a≤50, 11≤b≤18, 12≤c≤20, 6≤d≤10, 6≤ e ≦ 9, 1 ≦ f ≦ 20, and 0.01 ≦ g ≦ 0.5, and a + b + c + d + e + f + g = 100.

According to one embodiment of the invention, the complex may have a structure composed of a crystalline phase with dendritic morphology, and the size of the crystalline phase located within the amorphous matrix may be from several nanometers up to about 100 μm.

According to one embodiment of the present invention, the elements for forming the amorphous phase may be Ti, Zr, Be, Cu, Ni and the like.

According to one embodiment of the invention, the elements for the formation of the crystalline phase may comprise Ti, Zr, Nb, and / or Ta.

According to one embodiment of the invention, the atomic ratio Ti / Zr is less than 2.8, Ti / Be is less than 2.25, and Ti / (Cu + Ni) is less than 2.37 in the composition optimized for the formation of the amorphous structure. have.

According to one embodiment of the invention, about 5 at% <Nb <about 20 at% and about 2 at% <Ta <about 8 at% in a composition optimized for formation of the composite.

According to one embodiment of the present invention, the composite may have a mechanical strength of about 1600 MPa to about 2200 MPa and a density of about 6.1 g / cm ³ to about 8.0 g / cm ³ at room temperature.

According to one embodiment of the present invention, the casting condition, the complex at about 350 degrees Celsius, depending on the volume fraction on the bcc composition from about 2x10 ^-8 mol / msPa ^{0 .5} to about 6x10 ^-8 mol / high hydrogen of ^0.5 msPa Permeability values, which are greater than the hydrogen permeability values of pure palladium metals and palladium alloys prepared under the same conditions.

According to one embodiment of the invention, the composite may have a high corrosion resistance under a hydrogen environment. With H ₂ bubbling at about 80 degrees Celsius, the corrosion current density is lower than that of SUS-316L stainless steel or Ti-6Al-4V alloy. For example, 1M H ₂ SO ₄ and 2 ppm F as ring H ₂ bubble in degrees 80 degrees ^from the corrosion current density of the _{Ti 45 Zr 13 .1 Be 16 .4} Cu 8 .2 Ni 7 .3 Nb 10 is about 9.6 × 10 ⁻³ mA / cm ² , and under the same conditions, the corrosion current densities of SUS-316L and Ti-6Al-4V are about 1.5 mA / cm ² and about 0.9 mA / cm ^2, respectively.

According to one embodiment of the present invention, a method for producing a Ti-based bulk amorphous matrix composite includes vacuum suction casting technique and vacuum squeeze casting technique using a copper or iron mold of about 12 mm diameter.

According to one embodiment of the present invention, the method for producing a Ti-based bulk amorphous matrix composite comprises from about 0.6Tg to about about 1 hour or less for a relatively short time to release eventual residual stress. Heat treatment at a temperature of 0.8Tg. Tg is the glass transition temperature expressed by Calvin temperature here.

According to one embodiment of the present invention, the composite may be transformed into a complex shape in the supercooled liquid region under a neutral and vacuum environment. In the supercooled liquid region, 유리 T _x is the difference between the glass transition temperature T _g and the crystallization temperature T _x , ie ㅿ T _x = T _x -T _g . For example, a hydrogen film can be produced by the stamping technique at about 350 degrees Celsius to about 400 degrees Celsius by applying a small load for several minutes.

EXAMPLES Hereinafter, the present invention will be described in more detail with reference to Examples. However, the following Examples are merely examples of the present invention, but the present invention is not limited to the following Examples.

< Example  1>

Amorphous  Complex formation

The amorphous structure is composed of two kinds of alloys such as Ti-Cu and Ti-Ni, three kinds of alloys such as Ti-Cu-Zr, Ti-Zr-Ni, and Ti-Zr-Be, and four kinds of Ti-Zr-Cu-Ni. The alloy may be formed of a Ti alloy such as five alloys such as Ti-Zr-Be-Cu-Ni. For Ti-Zr-Be-Cu-Ni, the glass formation range can be quite extended and the maximum size reported so far is about 10 mm diameter. In addition, these alloys have a high strength on the order of about 1.6 GPa to about 2 GPa and can be modified with plastic deformation of up to about 10% under compression mode at room temperature. Mechanical properties can be improved by introducing small crystalline phases. Formation of the in situ crystalline phase in the amorphous phase can be accomplished by several methods. For example, a partial crystallization method by heat treatment of Tx or more for a short time, a phase separation method by immiscibility between two elements, or a phase partitioning method upon cooling. In this embodiment, Ti-Zr-Be-Cu-Ni alloy is applied.

Competitive phases in Ti-Zr-Be-Cu-Ni alloys include icosahedral quasicrystalline phase, TiZrCu ₂ or TiZrNi type Laves phase, α (Ti) and β ( Ti) phase. The icosahedral phase and the Laves phase are intermetallic phases and are generally avoided because of their fragile nature. The β (Ti) phase has a body centered cubic (bcc) structure that is softer than the hexagonal α (Ti) phase. The β (Ti) phase is common not only to Group 5 metals such as V, Nb and Ta, but also to Ti and Zr. In the phase diagrams shown in FIGS. 1A and 1B, the β (Ti) phase can be seen as having a high temperature phase throughout the composition. Although the α (Ti) phase is the most stable phase at room temperature, the formation of the α (Ti) phase in many alloys generally occurs slowly. Long heat treatment at low temperature is generally applied in order to surely control the volume fraction of the α phase. Although the α phase can form a Ti-Ta alloy having a wide range of compositions at low temperatures, Nb and Ta stabilize the β (Ti) phase. Due to the cooling rate of 1 K / s or more in vacuum casting technology, the phase most likely to occur during solidification is the β phase.

Thus, the addition of one or more Group 5 elements can promote the formation of the β phase by reducing the free energy of the β phase. 1C, 1D, and 1E, Nb and Ta have atomic radii of about 0.146 nm and about 0.147 nm, respectively, and [Kr] 4d4 5s1 (2.8.18.12.1 and electron-to-atomic ratios) to-atom ratio, e / a) = 5.4) and [Xe] 4f14 5d3 6s2 (2.8.18.32.11.2 and e / a = 5.5). Nb and Ta have an atomic radius similar to Ti, and can moderately promote glass formation by the confusion principle.

Another parameter is heat of mixing, and the heat mixing values for the major elements are shown in Table 1 below. Thermal mixing of Nb and Ta with major elements such as Ti, Zr, Cu is almost zero, and thermal mixing with Be and Ni has a negative value.

Although Nb and Ta are added to promote the formation of β (Ti, Zr) phases by the reduction of free energy, as shown in FIG. 1F, the physical, electrical and thermodynamic properties of these elements do not form glass in Ti-based amorphous metal Does not appear to seriously deteriorate and certainly forms a matrix with an amorphous structure. Table 1 shows the thermal mixing (kj / mol) values.

Ti Zr Be Cu Ni Ti - 0 -30 -9 -35 Zr 0 - -43 -23 -49 Be -30 -43 - 0 -4 Cu -9 -23 0 - 4 Ni -35 -49 -4 4 - Nb 2 4 -25 3 -30 Ta One 3 -24 2 -42

< Example  2>

Fabrication of Membrane Alloys

The bulk metallic glass and in situ BMG matrix composites produced in this embodiment can be formed by two different techniques, vacuum suction casting or vacuum squeeze casting.

Vacuum suction casting technology

Vacuum suction casting technology has been used to produce cylindrical rods. 2A is a view schematically showing a suction casting apparatus. The sample piece is placed over a water cooled copper mold with a small nozzle. The chamber is evacuated and under high purity argon gas atmosphere. The sample is then melted again and cast into a water-cooled copper mold with cylindrical pores 10 mm in diameter and 50 mm long or plate-shaped pores 12 mm wide, 2 mm thick and 50 mm long. . A film about 0.5 mm thick is cut and formed by wire cutting.

vacuum Squeeze  Casting technology

In situ BMG matrix composite samples can be prepared by the vacuum squeeze casting technique shown in FIG. 2B. In this technique, the ingot is melted first, and the upper die moves rapidly downward to pressurize the molten ingot on the hearth. The material of the upper die and stove can be copper, stainless steel, or the like. The size and shape of the sample can be determined by the pore size of the stove, and the sample can be cooled by a cooling system that ensures a fast cooling rate. By controlling the vertical movement of the upper die, the thickness of the sample can be adjusted. To prepare a sample having a composite structure, the upper die can be moved slowly so that the crystalline phase can grow.

< Example  3>

Ti - Zr - Be - Cu - Ni - Nb  Complex and Ti - Zr - Be - Cu - Ni - Ta  Structure of complex

FIG. 3A shows Ø10 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ and Ø7 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ prepared by vacuum casting ketchup. XRD graph of the alloy. Halo peaks are obtained when (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ and (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ are in the as-cast state. It is shown to be completely amorphous. Also DSC graphs of these alloys are shown in FIG. 3B. Ø10 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ and Ø8 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ bars have two exothermic peaks of about 425 degrees Celsius and about 525 degrees Celsius This corresponds to the transition from amorphous to metastable crystalline phase and from amorphous + metatable to Laves phase, respectively.

The structural thermal properties of the bulk alloy using (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb _{5 indicate} that Nb has little effect on the glass forming ability of the Ti-based alloy in this composition. Conversely, Ø10 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ The XRD graph of the alloy shows crystalline peaks, indicating that Ta significantly reduced the formation of amorphous structures by promoting the formation of the Laves crystalline phase, as shown in FIG. 3C.

The structure of the cast rod is confirmed by the SEM pictures shown in FIGS. 3D and 3E. Ø8 mm of FIG. 3E (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ The image of the rod shows the dendritic structure on β (Ti) in the amorphous matrix, Ø8 mm (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb _{5 of} FIG. The image of the bar does not represent any image. The complete amorphous structure of this alloy is confirmed by TEM observation.

A thermal feature of amorphous metals is the presence of supercooled liquid regions. In the supercooled liquid region, 유리 T _x is the difference between the glass transition temperature T _g and the crystallization temperature T _x , ie ㅿ T _x = T _x -T _g . The larger the supercooled liquid region, the easier the formation of the amorphous alloy. For such Ti-based alloys, Tg is difficult to detect and the supercooled liquid region can be difficult to define using calorimeter technology.

< Example  4>

Manufactured by vacuum suction casting technology Ti - Zr - Be - Cu - Ni - Nb Inshi Amorphous  Formation range of the metal matrix composite

In situ amorphous metal matrix composites (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ alloys produced by vacuum suction casting technology can be up to 10 mm in diameter. Nb containing alloys were tested in detail. Formation of the amorphous metal matrix composite was tested in the composition range shown in Figure 4a, the composition ratio is shown in Table 2 below. The Ti content can be approximately 38 to 50 at%, above 50 at% the glass forming ability rapidly decreases and the amorphous structure cannot be made of large sized material. This indicates that the Ti / Zr ratio value is between about 2.8 and 4.2 and the Ti / (Cu + Ni) ratio value is between about 2 and 3.8. Be also controls the glass forming ability. When the Be content is reduced, the formation of an amorphous matrix requires faster cooling rates. Thus, the Ti / Be ratio value is maintained between approximately 2.0 and 4.0. Table 2 below shows the composition (at%) of the in situ Ti-based amorphous metal matrix composite to which Nb was added.

Configuration Ti Zr Be Cu Ni Nb
A A1 42.75 15.2 19.0 9.5 8.55 5.0 A2 40.5 14.4 18.0 9.0 8.1 10.0 A3 38.25 13.6 17.0 8.5 7.65 15.0
B B1 45.0 14.55 18.18 9.09 8.18 5.0 B2 45.0 13.09 16.36 8.19 7.36 10.0 B3 45.0 11.64 14.54 7.28 6.54 15.0
C C1 38.25 15.055 18.82 9.405 8.47 10.0 C2 42.75 13.745 17.18 8.595 7.73 10.0 C3 47.25 12.435 15.54 7.785 6.99 10.0 C4 41.625 12.62 15.77 7.89 7.095 15.0
B M1 50.0 14.09 12.36 6.19 7.36 10.0 M2 47.5 14.0 14.0 8.2 6.3 10.0 M3 46.12 15.1 17.0 9.6 7.18 5.0

When exceeding 5 at%, the addition of Nb reduces the free energy on the bcc phase, thus promoting the nucleation of the crystalline phase in the liquid. As shown in Figures 4B and 4C, as the concentration of Nb increases, the volume fraction and size of the bcc crystalline phase also increases. This can be seen by the increase in peak intensity on bcc in the XRD graph shown in FIG. 4D.

The effect of Nb on the microstructure can be seen by the DSC graph shown in FIG. 4E. This can be seen by the reduction of the first and second exothermic peaks corresponding to the change from amorphous to metastable phase and the change from amorphous + metatable to crystalline phase, respectively. The first exothermic peak disappeared when Nb was 15 at%, indicating that the amorphous matrix directly transformed into the crystalline phase.

The β phase grows in a dendritic structure in the liquid, and this growth occurs along any crystallographic direction. However, the morphology of the β phase can be largely determined by the composition of the alloy. For example, FIGS. 4F and 4G show different morphologies with many small dendrites formed from numerous nucleation sites of Ti ₄₅ (Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) _45/55 Nb ₁₀ alloy. In the 4g Ti _41.6 Zr _12.6 Be _15.8 Cu _7.9 Ni _7.1 Nb ₁₅ alloy, the dendrites exhibited an uninterrupted morphology showing the growth of the β phase from one nucleation site. This difference in morphology has a great influence on the properties of the material. Table 3 below shows the compositional analysis of the matrix and dendrites by EDS (at%).

matrix A2 B1 B2 B3 C3 C4 Ti 41.5 52.6 48.0 47.5 49.4 48.8 Zr 25.5 17.8 19.7 18.2 18.2 19.0 Nb 9.5 8.2 6.5 7.6 9.0 6.7 Cu 12.0 10.5 14.1 13.9 11.8 13.2 Ni 11.5 10.9 11.8 12.8 11.6 12.3

Dendrite Ti 51.3 - 58.5 51.1 60.3 51.1 Zr 9.4 - 6.8 5.6 - - Nb 36.3 - 34.7 43.3 39.7 48.9 Cu - - - - - - Ni - - - - - -

< Example  5>

vacuum Squeeze  Manufactured by casting technology Ti - Zr - Be - Cu - Ni - Nb Inshi Amorphous  Formation of Metal Matrix Composites

Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ , (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ , (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₀ Nb ₁₀ , and (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₈₅ Nb ₁₅ alloys can be produced by vacuum squeeze casting technology in a plate shape of about 2-3 mm thick. XRD and DSC graphs for these samples are shown in FIGS. 5A and 5B. When the alloy compositions were identical, the structure of the alloy produced by the squeeze casting technique was found to be similar to the samples produced by the vacuum suction casting technique.

As shown in Figures 5c, 5d, 5e, 5f, the samples in the SEM photographs show a composite structure in the alloy containing Nb content greater than 10 at%.

< Example  6>

Manufactured by vacuum suction casting technology Ti - Zr - Be - Cu - Ni - Ta Inshi Amorphous  Formation of Metal Matrix Composites

(Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ alloy is produced in a plate shape with a thickness of about 1 mm by vacuum suction casting technique. As shown in the XRD graphs and SEM photographs shown in FIGS. 6A and 6B, respectively, an additional β (Ti) phase of 5 at% was formed in this alloy. The morphology of the dendritic β phase in Ta-containing in situ Ti-based BMG matrix composites is similar to that of Nb-containing in situ Ti-based BMG matrix composites. However, the addition of less Ta content than Nb leads to the formation of β (Ti) phases in the amorphous matrix more efficiently than Nb. For example, the volume of β phase in Ti and 5 at% Ta containing BMG matrix composites is similar to Ti and 10 at% Nb containing BMG matrix composites.

As shown in FIG. 6C, the amorphous properties of the matrix are confirmed by DSC analysis. The positions of the peaks as well as the heat of crystallization are similar to the Nb containing in situ Ti-based BMG matrix composites.

< Example  7>

Inshi Ti system BMG  Mechanical properties of the composite

Amorphous materials exhibit high microhardness values, high strength with Vickers microhardness greater than 300 kgf / mm ² (3 GPa), strengths near or above 2000 MPa (2 GPa) according to the composition ratios of Table 4. These characteristics are important factors for applications such as micro-gear, bipolar plate, etc., because they can reduce wear in the case of microgear, and the plate thickness is much thinner due to the high strength in the bipolar plate. This is because it can make a difference and can benefit from the weight and volume of the fuel cell.

The stress-strain curves of the Nb and Ta containing Ti based composites are shown in FIG. 7A. Compression tests were performed on samples with different diameters, but both Nb and Ta may exhibit some malleability at room temperature. Changes in plastic strain values from one alloy to another can be explained by differences in the volume fraction, size, and properties of the crystalline dendritic phase. Table 4 shows the microhardness values of various Ti-based BMG or BMG composites.

Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Ta ₅ Hv 366 ± 22 344 ± 24 414 ± 42 (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₀ Nb ₁₀ (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₈₅ Nb ₁₅ Hv 344 ± 24 375 ± 37 366 ± 14

Nb-containing in situ Ti-based BMG composite samples are prepared using a stainless steel mold instead of a Cu mold to reduce cooling and increase the volume fraction on the bcc. As shown in FIG. 7B, a change in sample preparation can change the mechanical behavior under the compression mode of the alloys. As shown in FIG. 7A, plastic deformation values of about 5-8% were obtained instead of about 2-3% when the volume fraction on bcc was increased.

For application as a material for hydrogen separation and purification membranes, in situ Ti-BMG composites preferably have a rather high strength at membrane application temperatures of about 300-400 degrees Celsius. The properties of these materials are present in the supercooled liquid region, for example, ㅿ Tx between the glass transition temperature Tg and the crystallization temperature Tx. In this temperature range, BMG can be transformed into superplastic behavior with high strain rate sensitivity.

7C, 7D and 7E are in situ containing 5, 10 and 15 at% Nb measured at 360 degrees Celsius at strain rates of 10 ⁻⁵ s ⁻¹ , 10 ⁻⁴ s ⁻¹ and 10 ⁻³ s ⁻¹ , respectively The stress-strain behavior of Ti-based BMG complexes is shown. Such composites, unlike tests measured at room temperature as shown in FIGS. 7A and 7B, can deform into large plastic deformations at high temperatures while maintaining high strength.

Strain rate sensitivity can be determined by variation of flow stress with strain rate. As shown in Fig 7f, at 5% Nb-containing alloy as well as the Nb-free alloys are both integral (monolithic) shows a high strain rate sensitivity as a BMG alloy, slope m represents the strain rate sensitivity is 10 ⁴ - 10 ^{- It} is approximately 0.5 (m = (δlogσ / (δlogε) near the strain rate of ⁵ s ⁻¹ . The structure of the in situ BMG complex consists of the bcc crystalline phase in the amorphous matrix, and the gradient value decreases with increasing Nb content. This behavior is due to the appearance of β phase in the amorphous matrix, which has a low strain rate sensitivity compared to the amorphous structure.

< Example  8>

Inshi Ti - Zr - Be - Cu - Ni + Nb BMG  Hydrogen Permeability Test of Composites

The hydrogen permeability test was performed on a sample of about 12 mm in diameter and about 0.5 mm in thickness. These sample films are made from plates formed by different methods of vacuum squeeze casting technology or vacuum suction casting technology. The test was carried out between hydrogen pressure of 1.1 to 3 bar, 350 to 375 degrees Celsius.

With respect to samples made by vacuum squeeze casting technology, FIG. 8A shows that the hydrogen flux increases at 350 degrees Celsius as the Nb content increases, indicating that Nb free and 5 at% Nb containing alloys are not composites. It is due to the appearance of bcc crystalline dendritic structures in in situ BMG complexes because they are integral amorphous metals. However, the difference in properties between Nb free and 5 at% Nb containing alloys suggests a positive effect due to Nb addition in hydrogen permeation.

Similar behavior is observed in samples made by vacuum suction casting techniques. The hydrogen permeation properties of the in situ Ti-based BMG composites are superior to the hydrogen permeation properties of pure palladium, as shown in FIG. 8B, and FIG. 8B shows the change in flux of the material. . In FIG. 8B the amount of change is almost linear, indicating that the change in hydrogen concentration follows the Sievert law. 8B also shows that the hydrogen flux increases as the temperature increases.

These properties are also shown in Figures 8c and 8d, which show the change in hydrogen permeability with the inverse of the temperature, wherein the hydrogen permeability value of the composite containing 10b or more of Nb is the hydrogen permeability value of palladium and pure Nb hydrogen. Indicates that the transmission is between values.

8C and 8D, the hydrogen permeability of the composite measured at 350 degrees Celsius and 375 degrees Celsius is compared with the value of the Ni—Nb—Zr based monolithic metallic glass. The hydrogen permeability of the composite is approximately 2.5 times greater than the hydrogen permeability of the palladium alloy and approximately 3.5 times greater than the Ni—Nb—Zr based integral amorphous metal.

The hydrogen permeability characteristics of the samples produced by the vacuum suction casting technique appear to be slightly better than the hydrogen permeability characteristics of the samples prepared by the vacuum squeeze casting technique. This is because the difference in cooling rate causes a difference in volume fraction and morphology of the bcc dendritic phase, which causes a difference in structure.

Inshi Ti - Zr - Be - Cu - Ni + Ta BMG  Hydrogen Permeability Test of Composites

The hydrogen permeability test was performed on a sample of ₉₅ Ta ₅ (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) prepared by vacuum squeeze casting technique as described in Example 6. As shown in FIGS. 8E and 8F, the hydrogen permeability of the 5 at% Ta containing composite is similar to that of the 10 at% Nb containing composite.

The hydrogen permeability value is high because the volume fraction of the dendrites formed by the addition of 5 at% Ta is high.

< Example  9>

Since the constituent elements used in these alloys have good corrosion resistance, the prepared in situ BMG composites also exhibit excellent corrosion resistance. Ti, Zr, and Nb have excellent corrosion resistance in H ₂ SO ₄ solution, Cu is a noble element, Ni also has good corrosion resistance, and Be forms BeO with excellent corrosion resistance.

Corrosion experiments were carried out at 1 MH ₂ SO ₄ and 2 ppm F ⁻ with hydrogen bubbling at 80 degrees Celsius to investigate the resistance of the composite in a hydrogen environment, where a potentio-dynamic graph is shown in FIG. 9A. Was shown as.

The potentiodynamic graph of all composites was found to be lower than the stainless steel and Ti-6Al-4V alloys tested in the same environment, indicating that the in situ Ti-based BMG composites had excellent corrosion resistance. In addition, unlike stainless steel and Ti-6Al-4V alloys, the polarity graph for Ti-based BMG composites does not exhibit active-passive transition behavior.

The behavior of Ti-6Al-4V alloys shows a wide passivation plateau followed by a passive dissolution without significant increase in current. Corrosion parameters are shown in Table 5 below.

Under hydrogen bubbling conditions, 10 at% containing alloys exhibit a low passivation current density when compared to alloys of other compositions. As shown in Table 6 below, the formation of (Ti, Zr, Nb) -rich β phase in a 15 at% Nb-containing composite increases the concentration of copper in the matrix, which tends to reduce corrosion resistance.

As shown in FIG. 9B, the change in corrosion current density and protection current density indicates that the corrosion resistance can be improved with a suitable alloy design, although the addition of Nb can result in the formation of inheterogeneity. Table 5 below shows the values of corrosion current density and potential corrosion at 1 M H ₂ SO ₄ and 2 ppm F ⁻ with hydrogen or air bubbling at 80 degrees Celsius.

H ₂ bubbling Air bubbling I _corr (A / cm ² ) E _corr (V) I _corr (A / cm ² ) E _corr (V) Ti-6Al-4V 9 x 10 ^-4 -0.76 7.2 x 10 ^-4 -0.78 Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ 3.01 x 10 ^-5 -0.23 4.73 x 10 ^-5 -0.13 (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ 2.03 x 10 ^-5 -0.23 1.34 x 10 ^-5 -0.17 (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₀ Nb ₁₀ 9.14 x 10 ^-6 -0.19 1.03 x 10 ^-5 -0.13 (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₈₅ Nb ₁₅ 8.11 x 10 ^-6 -0.23 1.39 x 10 ^-5 -0.17 _{Ti 45 Zr 14 .55 Be 18 .18} Cu 9 .09 Ni 8 .18 Nb 5 2.53 x 10 ^-5 -0.26 1.62 x 10 ^-5 -0.17 _{Ti 45 Zr 13 .09 Be 16 .36} Cu 8 .19 Ni 7 .36 Nb 10 2.29 x 10 ^-5 -0.25 1.83 x 10 ^-5 -0.21 _{Ti 45 Zr 11 .64 Be 14 .54} Cu 7 .28 Ni 5 .54 Nb 15 2.13 x 10 ^-5 -0.27 1.47 x 10 ^-5 -0.26 (Ti ₄₀ Zr ₂₉ Be ₁₆ Cu ₈ Ni ₇ ) ₉₅ Nb ₅ - - 4.56 x 10 ^-6 -0.20 (Ti ₄₀ Zr ₂₉ Be ₁₆ Cu ₈ Ni ₇ ) ₉₇ Ta ₃ - - 3.42 x 10 ^-6 -0.28 (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₈₀ Nb ₁₈ Ta ₂ 3.92 x 10 ^-6 -0.27 3.03 x 10 ^-6 -0.19

Potentio-static tests were performed at a potential constant value of -0.1 V for 2 hours. The change in current density of the in situ BMG is shown in FIG. 9C and compared with the change in current density of Ti-6Al-4V. After two hours of testing, the slope of the current density versus time appears to decrease as the Nb content increases from 0 to 10 at% with the absolute value of the 10 at% Nb-containing composite lower than the alloy of the other composition. However, the slope of the 15 at% Nb containing composite is too large, indicating that the corrosion resistance is lost under a hydrogen environment.

Nb-free integral Ti-based BMG and Nb-containing Ti-based BMG composites were immersed in 1M H ₂ SO ₄ at 80 degrees Celsius for 7 days. The results shown in Table 6 below show that the weight loss of Ti-based BMG composites recorded after 7 days is significantly lower than the weight loss of crystalline Ti-6Al-4V, which shows that the corrosion resistance of the manufactured BAMC is very good. . In addition, the best corrosion resistance is seen in composites containing 10 at% Nb. Table 6 below shows the weight loss after soaking in 1M H ₂ SO ₄ at 80 degrees for 7 days, the weight loss being measured using a fine balance and calculated from the amount of elements dissolved in the solution.

Alloy % of weight loss
Immersion
time Weight loss method Atomic Absorption Spectroscopy (AAS) method Ti-6Al-4V (crystal) 34.12 37.32 2 days 3 hrs 27 min Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ 5.39 4.53 7 days (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₅ Nb ₅ 2.15 1.82 7 days (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₉₀ Nb ₁₀ 1.59 1.33 7 days (Ti ₄₅ Zr ₁₆ Be ₂₀ Cu ₁₀ Ni ₉ ) ₈₅ Nb ₁₅ 1.82 1.73 7 days (Ti ₄₀ Zr ₂₉ Be ₁₆ Cu ₈ Ni ₇ ) ₉₅ Nb ₅ 1.72 7 days (Ti ₄₀ Zr ₂₉ Be ₁₆ Cu ₈ Ni ₇ ) ₉₇ Ta ₃ 0.78 0.88 7 days

Although the preferred embodiments of the present invention have been described in detail above, the scope of the present invention is not limited thereto, and various modifications and improvements of those skilled in the art using the basic concepts of the present invention defined in the following claims are also provided. It belongs to the scope of rights.

Claims (12)

A Ti-based in-situ bulk amorphous matrix composite having a composition (at%) represented by Formula 1 below:
[Chemical Formula 1]
Ti _a Zr _b Be _c Cu _d Ni _e M _f I _g
In Formula 1, M is at least one of Nb or Ta, I is impurity, 38≤a≤50, 11≤b≤18, 12≤c≤20, 6≤d≤10, 6≤e≤9, 1 ≤ f ≤ 20, and 0.01 ≤ g ≤ 0.5, and a + b + c + d + e + f + g = 100. In claim 1,
The complex has a structure consisting of a crystalline phase having a dendritic morphology (dendritic morphology), the size of the crystalline phase located in the amorphous matrix is in situ bulk amorphous matrix composite of 0.01 μm to 100 μm. 3. The method of claim 2,
An element for forming the amorphous matrix is a Ti-based in-situ bulk amorphous matrix composite comprising Ti, Zr, Be, Cu, and Ni. 3. The method of claim 2,
The element for forming the crystalline phase is a Ti-based in-situ bulk amorphous matrix composite comprising at least one of Ti, Zr, and Nb or Ta. 3. The method of claim 2,
Ti-based in situ bulk amorphous matrix composite having an atomic ratio Ti / Zr of less than 2.8, Ti / Be of less than 2.25, and Ti / (Cu + Ni) of less than 2.37 in the composition for forming the amorphous matrix. The method of claim 5,
Ti-based in-situ bulk amorphous matrix composite of 5 at% <Nb <20 at% and 2 at% <Ta <8 at% in the composition. In claim 1,
The composite is a Ti-based in-situ bulk amorphous matrix composite having a mechanical strength of 1600 MPa to 2200 MPa and a density of 6.1 g / cm ³ to 8.0 g / cm ³ at room temperature. In claim 1,
The Ti-based composite in-situ bulk amorphous matrix composites having a hydrogen permeability value at 350 degrees Celsius according to the composition of 2x10 ^-8 mol / msPa ^{0 .5} to 6x10 ^-8 mol / msPa ^{0 .5.} In claim 1,
The composite is a Ti-based in-situ bulk amorphous matrix composite having a lower corrosion current density than stainless steel under a hydrogen environment. Vacuum suction casting technique and vacuum squeeze casting technique,
To prepare a Ti-based in-situ bulk amorphous matrix composite having a composition represented by the following formula (at%):
[Chemical Formula 1]
Ti _a Zr _b Be _c Cu _d Ni _e M _f I _g
In Formula 1, M is at least one of Nb or Ta, I is impurity, 38≤a≤50, 11≤b≤18, 12≤c≤20, 6≤d≤10, 6≤e≤9, 1 ≤ f ≤ 20, and 0.01 ≤ g ≤ 0.5, and a + b + c + d + e + f + g = 100. 11. The method of claim 10,
Ti-based further comprising a heat treatment at a temperature of 0.6Tg to 0.8Tg (Tg is the glass transition temperature expressed in Calvin temperature) for a predetermined time of less than 1 hour to release the eventual residual stress (release) A method of making an in situ bulk amorphous matrix composite. 11. The method of claim 10,
The complex is transformed into a complex shape in the supercooled liquid region under a neutral and vacuum environment, in which ㅿ T _x is the crystallization temperature and glass transition temperature T _g in the supercooled liquid region. A method of preparing a Ti-based in-situ bulk amorphous matrix composite which is a difference of temperature T _x .

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