EP0905268A1

EP0905268A1 - High-strength amorphous alloy and process for preparing the same

Info

Publication number: EP0905268A1
Application number: EP98111771A
Authority: EP
Inventors: Akihisa Inoue; Tao Zhang; Hidenobu Nagahama
Original assignee: YKK Corp
Current assignee: YKK Corp
Priority date: 1997-08-29
Filing date: 1998-06-25
Publication date: 1999-03-31
Also published as: JPH1171661A

Abstract

A high-strength amorphous alloy represented by the general formula: X_aM_bAl_cT_d (wherein X is at least one element selected between Zr and Hf; M is at least one element selected from the group consisting of Ni, Cu, Fe, Co and Mn; T is at least one element having a negative enthalpy of mixing with at least one of the above-mentioned X, M and Al; and a, b, c and d are atomic percentages, provided that 25 ≦ a ≦ 85, 5 ≦ b ≦ 70, 0 < c ≦ 35 and 0 < d ≦ 15) and having a structure comprising at least an amorphous phase. The element T may be at least one element selected from the group consisting of Ru, Os, Rh, Ir, Pd, Pt, V, Nb, Ta, Cr, Mo, W, Au, Ga, Ge, Re, Si, Sn and Ti. The amorphous alloy has high hardness and strength, excellent ductility, high corrosion resistance, and excellent workability and produced by preparing an amorphous alloy having the above composition and heat-treating in the specified temperature range to decompose the amorphous phase into a mixed phase structure of an amorphous phase and a microcrystalline phase.

Description

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to an amorphous alloy having high hardness and strength, excellent ductility, high corrosion resistance, and excellent workability, and a process for preparing the same.

2. Description of the Prior Art

Conventional Zr-based alloys having specified alloy compositions causes glass transition before crystallization, have a wide supercooled liquid region, and have a high capability of forming an amorphous phase. Since these alloys have such a high amorphizing capability, they become amorphous not only by any method wherein a high cooling rate can be secured like a liquid quenching method, but also by any ordinary casting method wherein the cooling rate is slow like a copper mold casting method, whereby tough bulk amorphous alloys can be prepared. When, however, a quenched tough thin strip formed by, for example, the liquid quenching method is heated at a temperature around the crystallization temperature thereof to precipitate crystals, the toughness thereof is deteriorated so that it can hardly be subjected to 180° contact bending. On the other hand, according to the copper mold casting method, a good amorphous bulk can be formed when cooled at a given or higher cooling rate, while the toughness thereof is deteriorated when the cooling rate is lowered to precipitate crystals.

SUMMARY OF THE INVENTION

The present invention aims at providing a high-strength amorphous alloy while solving the problem of deterioration of toughness either when a formed quenched tough thin strip or bulk material is heat-treated to precipitate crystals or when the cooling rate is lowered in the mold casting method to precipitate crystals.
The present invention provides a high-strength amorphous alloy represented by the general formula: X_aM_bAl_cT_d (wherein X is at least one element selected between Zr and Hf; M is at least one element selected from the group consisting of Ni, Cu, Fe, Co and Mn; T is at least one element having a negative enthalpy of mixing with at least one of the above-mentioned X, M and Al; and a, b, c and d are atomic percentages, provided that 25 ≦ a ≦ 85, 5 ≦ b ≦ 70, 0 < c ≦ 35 and 0 < d ≦ 15) and having a structure comprising at least an amorphous phase.
The above-mentioned element T is at least one element selected from the group consisting of Ru, Os, Rh, Ir, Pd, Pt, V, Nb, Ta, Cr, Mo, W, Au, Ga, Ge, Re, Si, Sn and Ti, among which Pd, Pt and Au are especially effective.
The addition of such an element T can bring about a change in the bonding of the constituent elements of the resulting amorphous alloy so as to allow it to attain a high strength without deterioration of toughness. Further, the structure of the alloy of the present invention is a mixed phase comprising an amorphous phase and a microcrystalline phase. The formation of the mixed phase structure provides excellent mechanical strength and ductility. When particular consideration is given to ductility, the amorphous phase preferably accounts for at least 50% in terms of volume fraction.
The present invention also provides a process for preparing a high-strength amorphous alloy, comprising preparing an amorphous alloy having a composition represented by the aforementioned general formula and containing at least an amorphous phase, and heat-treating the alloy in the temperature range from the glass transition temperature Tg thereof to the first exothermic reaction-starting temperature (Tx₁: crystallization temperature) thereof to decompose the amorphous phase into a mixed phase structure consisting of an amorphous phase and a microcrystalline phase.

BRIEF DESCRIPTION OF THE DRAWINGS

Fig. 1 is a graph showing the results of differential scanning calorimetric analyses of the alloys of the present invention and the alloy of Comparative Example.
Fig. 2 is a graph showing data on the X-ray diffraction analysis when the alloys of the present invention and the alloy of Comparative Example were subjected to a predetermined heat treatment.
Figs. 3A and 3B are the TEM and electron diffraction photographs of the alloy of Comparative Example, while Figs. 3C and 3D are the TEM and electron diffraction photographs of the alloy of the present invention.
Fig. 4 is a graph showing the mechanical properties of the alloys of the present invention.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

The above-mentioned amorphous alloy can be prepared by quenching a molten alloy having the above-mentioned composition according to a liquid quenching method such as a single roller melt-spinning method, a twin roller melt-spinning method, an in-rotating-water melt-spinning method, a high-pressure gas atomizing method, or a spray method, by rapidly cooling it according to sputtering, or by slowly cooling it according to a mold casting method.
The amorphous alloy thus obtained is heat-treated. When, however, it is heat-treated below Tg, it cannot easily be decomposed into a microcrystalline phase. On the other hand, when the heat-treating temperature exceeds Tx₁, the crystalline phase formed by decomposition cannot be inhibited from coarsening. Accordingly, the structure can be stabilized by effecting the heat-treating in the range of Tg to Tx₁.
The heating time may be 1 to 120 minutes. When it is shorter than 1 minute, no effect of heat-treating can be expected. When it exceeds 120 minutes, the crystalline phase is coarsened. This phenomenon is notably observed at a heat-treating temperature close to Tx₁.

Example 1

A mother alloy having the following composition: Zr₆₀Cu_30-xAl₁₀Pd_x (wherein x = 0, 5 or 10) (wherein the subscript refers to atomic %) was melted in an arc melting furnace, and then formed into a thin strip (thickness: 20 µm, width: 1.5 mm) with a single-roll liquid quenching unit (melt spinning unit) generally used. In this step, a roll made of copper and having a diameter of 200 mm was used at a number of revolutions of 4,000 rpm in an Ar atmosphere of not higher than 10^-3 Torr. The case where x = 5 or 10 corresponds to Example of the present invention, while the case where x = 0 corresponds to Comparative Example.
The resulting thin strip of the amorphous single-phase alloy was examined with a differential scanning calorimeter (DSC). In Fig. 1, the thermal properties of Zr₆₀Cu₃₀Al₁₀ (Comparative Example) are denoted by (a), while the thermal properties of Zr₆₀Cu₂₀Al₁₀Pd₁₀ and Zr₆₀Cu₂₅Al₁₀Pd₅ (Examples) are denoted by (b).
The glass transition temperature (Tg) and crystallization temperatures (Tx, Tx₁ and Tx₂) of each of the alloys (a) or (b) were as shown in the Figure. In this Figure, the supercooled liquid region (ΔT) is a region falling between the glass transition temperature (Tg) and each crystallization temperature (Tx, Tx₁ or Tx₂), a region where the control of temperature during working and the control of working time can be comparatively easily done, and a region serving as one of the criteria of whether or not working can be easily done. The workability can be judged by the temperature width (ΔT = Tx - Tg) corresponding to this region. More specifically, when ΔT is large, the working temperature width can be widened and also the working time can be lengthened, whereby working can be comparatively easily done.
A description will now be made of the method of determining Tg and Tx in the present invention. The Tg refers to a temperature at a point of intersection of the extrapolated base line with the rising portion of the differential scanning calorimetric curve in a region of the curve where an endothermic reaction occurs, while the Tx refers to a temperature found in the same manner in a region where an exothermic reaction occurs the other way around.
It is understood from Fig. 1 that the alloy of the present invention is an alloy having a wide supercooled liquid region over 50 K in width as compared with the alloy of Comparative Example. It is also understood that the alloy of the present invention is an alloy having two exothermic peaks (Tx₁, Tx₂).
Fig. 2 shows data on the X-ray diffraction analysis when the alloys (b) of Example as well as the alloy (a) of Comparative Example were heat-treated at a predetermined temperature for a predetermined period of time. It is understood that the alloy of Comparative Example underwent a considerable progress of crystallization by heat-treating in the supercooled liquid region at 705 K for 20 minutes. By contrast, it can be confirmed that the alloys of Example underwent no substantial change at 726 K in the supercooled liquid region for 30 minutes or 60 minutes, with a broad diffraction pattern peculiar to the amorphous phase. It can also be confirmed that the inventive alloys underwent a considerable progress of crystallization when heat-treated at 808 K above the supercooled liquid region, but still had a broad diffraction pattern as compared with the alloy of Comparative Example.
It is also understandable from the foregoing results that the alloy of the present invention is excellent in thermal stability and especially excellent in the properties in the supercooled liquid region Tg - Tx₁.
Figs. 3A and 3B are the microstructural photographs (TEM, electron diffraction) of the alloy of Comparative Example when it was heat-treated in the supercooled liquid region for 40 minutes, from which it is understood that the crystal grains thereof were considerably coarsened as compared with those of Example of the present invention, which will be described later. When this result is considered together with Fig. 2, it is conceivable that Zr₂Cu and Zr₂Al as crystal phases were coarsened.
By contrast, it is understood that a microcrystalline phase existed in a dispersed state despite heat-treating in the supercooled liquid region for 60 minutes in Example of the present invention [TEM of Fig. 3C]. When this result is considered together with Fig. 2, it is conceivable that the structure was such that a microcrystalline phase of Zr₂Cu dispersed in an amorphous phase. A halo pattern peculiar to an amorphous alloy can also be confirmed from the electron diffraction photograph thereof (Fig. 3D).
The alloys of Example were examined with respect to mechanical properties as against the volume fraction of the crystalline phase present in the matrix while varying the heat-treating time. The heating temperature refers to a temperature in the supercooled liquid region. As shown in Fig. 4, it is understandable that the mechanical properties thereof such as tensile strength, hardness and Young's modulus were improved in keeping with an increase in the volume of the microcrystalline phase dispersed in the amorphous phase. In this Figure, data at a volume fraction of 0 correspond to those of the alloys not heat-treated. Further, as a result of the 180° contact bending test for every sample in Fig. 4, it was found out that all materials were capable of contact bending and were endowed with an excellent ductility.
The alloy of the present invention is a material endowed not only with an excellent thermal stability and excellent mechanical properties but also with an excellent ductility. Further, according to the process of the present invention, a material stabilized in structure and endowed with the foregoing properties can be prepared with proper control.

Claims

A high-strength amorphous alloy represented by the general formula: X_aM_bAl_cT_d (wherein X is at least one element selected between Zr and Hf; M is at least one element selected from the group consisting of Ni, Cu, Fe, Co and Mn; T is at least one element having a negative enthalpy of mixing with at least one of the above-mentioned X, M and Al; and a, b, c and d are atomic percentages, provided that 25 ≦ a ≦ 85, 5 ≦ b ≦ 70, 0 < c ≦ 35 and 0 < d ≦ 15) and having a structure comprising at least an amorphous phase.
A high-strength amorphous alloy as claimed in claim 1, wherein said element T is at least one element selected from the group consisting of Ru, Os, Rh, Ir, Pd, Pt, V, Nb, Ta, Cr, Mo, W, Au, Ga, Ge, Re, Si, Sn and Ti.
A high-strength amorphous alloy as claimed in claim 1 or 2, wherein said structure is a mixed phase consisting of said amorphous phase and a microcrystalline phase.
A process for preparing a high-strength amorphous alloy, comprising preparing an amorphous alloy having a composition represented by the general formula: X_aM_bAl_cT_d (wherein X is at least one element selected between Zr and Hf; M is at least one element selected from the group consisting of Ni, Cu, Fe, Co and Mn; T is at least one element having a negative enthalpy of mixing with at least one of the above-mentioned X, M and Al; and a, b, c and d are atomic percentages, provided that 25 ≦ a ≦ 85, 5 ≦ b ≦ 70, 0 < c ≦ 35 and 0 < d ≦ 15) and containing at least an amorphous phase, and heat-treating said alloy in the temperature range from the glass transition temperature Tg thereof to the first exothermic reaction-starting temperature (Tx₁: crystallization temperature) thereof to decompose said amorphous phase into a mixed phase structure consisting of an amorphous phase and a microcrystalline phase.
A process for preparing a high-strength amorphous alloy as claimed in claim 4, wherein the heat-treating is effected in said temperature range for 1 to 120 minutes.
A process for preparing a high-strength amorphous alloy as claimed in claim 4 or 5, wherein said alloy containing an amorphous phase is an alloy consisting of an amorphous single phase.