High strength and anti-corrosive aluminum-based alloy

The present invention provides a high strength and anti-corrosive aluminum-based alloy essentially consisting of an amorphous structure or a multiphase amorphous/fine crystalline structure, which is represented by the general formula Al.sub.x M.sub.y R.sub.z. In this formula, M represents at least one metal element selected from the group consisting of Ti, V, Cr, Mn, Fe, Co, Cu, Zr, Nb, Mo and Ni, and R represents at least one element or mixture selected from the group consisting of Y, Ce, La, Nd and Mm (misch metal). Additionally, in the formula, x, y and z represent the composition ratio, and are atomic percentages satisfying the relationships of x+y+z=100, 64.5.ltoreq.x.ltoreq.95, 5.ltoreq.y.ltoreq.35, and 0<z.ltoreq.0.4.

BACKGROUND OF THE INVENTION 
1. Field of the Invention 
The present invention relates to an aluminum-based alloy for use in a wide 
range of applications such as in aircraft, vehicles and ships, as well as, 
in the structural material for the engine portions thereof. In addition, 
the present invention may be employed as sash, roofing material and 
exterior material for use in construction, or as material for use in sea 
water equipment, nuclear reactors, and the like. 
2. Description of Related Art 
As prior art aluminum-based alloys, alloys incorporating various components 
such as Al--Cu, Al--Si, Al--Mg, Al--Cu--Si, Al--Cu--Mg, and Al--Zn--Mg are 
known. In all of the aforementioned, superior anti-corrosive properties 
are obtained at a light weight, and thus the aforementioned alloys are 
being widely used as structural material for machines in vehicles, ships 
and aircraft, in addition to being employed as sash, roofing material, 
exterior material for use in construction, structural material for use in 
LNG tanks, and the like. 
However, the prior art aluminum-based alloys generally exhibit 
disadvantages such as a low hardness and poor heat resistance when 
compared to material incorporating Fe. In addition, although some 
materials have incorporated elements such as Cu, Mg and Zn for increased 
hardness, disadvantages remain such as low anti-corrosive properties. 
On the other hand, recently, experiments are being conducted in which the 
compositions of aluminum-based alloys are being refined by means of 
performing quench solidification from a liquid-melt state resulting in the 
production of superior mechanical strength and anti-corrosive properties. 
In Japanese Patent Application First Publication No. 1-275732, an 
aluminum-based alloy is disclosed which can be utilized as material with a 
high hardness, high strength, high electrical resistance, anti-abrasion 
properties, or as soldering material. In addition, the disclosed 
aluminum-based alloy has a superior heat resistance, and may undergo 
extruding or press processing by utilizing the superplastic phenomenon 
observed near liquid crystallization temperatures. This aluminum-based 
alloy comprises a composition AlM*X with a special composition ratio 
(wherein M* signifies an element such as V, Cr, Mn, Fe, Co, Ni, Cu, Zr and 
the like, and X represents a rare earth element such as La, Ce, Sm and Nd, 
or an element such as Y, Nb, Ta, Mm (misch metal) and the like), and has 
an amorphous or a combined amorphous/fine crystalline structure. 
However, this aluminum-based alloy is disadvantageous in that high costs 
result from the incorporation of large amounts of expensive rare earth 
elements and/or metal elements with a high activity such as Y. In addition 
to the aforementioned use of expensive raw materials, problems also arise 
such as increased consumption and labor costs due to the large scale of 
the manufacturing facilities required to treat materials with high 
activities. Furthermore, the aforementioned aluminum-based alloy tends to 
display insufficient resistance to oxidation and corrosion. 
SUMMARY OF THE INVENTION 
It is an object of the present invention to provide an aluminum-based 
alloy, possessing superior strength and anti-corrosive properties, which 
comprises a composition in which the incorporated amount of high activity 
elements such as Y or expensive elements such as rare earth elements is 
restricted to a small amount, or in which such elements are not 
incorporated at all, thereby effectively reducing the cost, as well as, 
the activity described in the aforementioned. 
In order to solve the aforementioned problems, the first aspect of the 
present invention provides an aluminum-based alloy, essentially consisting 
of an amorphous structure or a multiphase amorphous/fine crystalline 
structure, represented by the general formula Al.sub.x M.sub.y R.sub.z 
(wherein M is at least one metal element selected from the group 
consisting of Ti, V, Cr, Mn, Fe, Co, Cu, Zr, Nb, Mo and Ni, and R is at 
least one element or mixture selected from the group consisting of Y, Ce, 
La, Nd and Mm (misch metal)). In the formula, x, y and z represent the 
composition ratio, and are atomic percentages satisfying the relationships 
of x+y+z=100, 64.5.ltoreq.x.ltoreq.95, 5&lt;y .ltoreq.35, and 0&lt;z.ltoreq.0.4. 
The second aspect of the present invention provides an aluminum-based 
alloy, essentially consisting of an amorphous structure or a multiphase 
amorphous/fine crystalline structure, represented by the general formula 
Al.sub.x Ni.sub.y M' z (wherein M' is at least one metal element selected 
from the group consisting of Ti, V, Mn, Fe, Co, Cu and Zr). In the 
formula, x, y and z represent the composition ratio, and are atomic 
percentages satisfying the relationships of x+y+z=100, 
50.ltoreq.x.ltoreq.95, 0.5.ltoreq.y .ltoreq.35, and 
0.5.ltoreq.z.ltoreq.20. 
According to the third aspect of the present invention, the fine 
crystalline component of the multiphase structure described in the 
aforementioned first and second aspects comprises at least one phase 
selected from the group consisting of an aluminum phase, a stable or 
metastable intermetallic compound phase, and a metal solid solution 
comprising an aluminum matrix. The individual crystal diameter of this 
fine crystalline component is approximately 30 to 50 nm. 
The fourth aspect of the present invention provides an aluminum-based alloy 
represented by the general formula Al.sub.x Co.sub.y M".sub.z (wherein 
M"is at least one metal element selected from the group consisting of Mn, 
Fe and Cu). In the formula, x, y and z represent the composition ratio, 
and are atomic percentages satisfying the relationships of x+y+z=100, 
50.ltoreq.x.ltoreq.95, 0.5.ltoreq.y.ltoreq.35, and 0.5.ltoreq.z.ltoreq.20. 
The fifth aspect of the present invention provides an aluminum-based alloy 
represented by the general formula Al.sub.a Fe.sub.b L.sub.c (wherein L is 
at least one metal element selected from the group consisting of Mn and 
Cu). In the formula, a, b and c represent the composition ratio, and are 
atomic percentages satisfying the relationships of a+b+c=100, 
50.ltoreq.a.ltoreq.95, 0.5.ltoreq.b.ltoreq.35, and 0.5.ltoreq.c.ltoreq.20. 
The sixth aspect of the present invention substitutes Ti or Zr in place of 
element M"or L, in an amount corresponding to one-half or less of the 
atomic percentage of M" or L. 
In the aforementioned aluminium-based alloy according to the present 
invention represented by the formula Al.sub.x M.sub.y R.sub.z, the atomic 
percentages of Al, element M, and element R are restricted to 64.5-95%, 
5-35% and 0-0.4%, respectively. This is due to the fact that when the 
composition of any of the aforementioned elements fall outside these 
specified ranges, it becomes difficult to form an amorphous component, as 
well as a supersaturated solid solution in which the amount of solute 
exceeds the critical solid solubility; this, in turn, results in the 
objective of the present invention, an aluminum-based alloy having an 
amorphous structure, an amorphous/fine crystalline complex structure or a 
fine crystalline structure, being unobtainable using an industrial 
quenching process incorporating a liquid quenching method and the like. 
In addition, when diverging from the aforementioned composition ranges, it 
becomes difficult to obtain an amorphous phase for use in producing the 
fine crystalline complex structure, through crystallization of the 
amorphous phase produced by the quenching method using an appropriate 
heating process, or temperature control of a powder molding process which 
utilizes conventional powder metallurgy technology. 
Element M, which represents one or more metal elements selected from the 
group consisting of Ti, V, Cr, Mn, Fe, Co, Cu, Zr, Nb, Mo and Ni, coexists 
with R and improves the amorphous forming properties, as well as, raising 
the crystallization temperature of the amorphous phase. Most importantly, 
this element markedly improves the hardness and strength of the amorphous 
phase. 
As well, under the fine crystal manufacturing conditions, these elements 
also stabilize the fine crystalline phase, form stable or metastable 
intermetallic compounds with aluminum or other additional elements, 
disperse uniformly in the aluminum matrix (.alpha.-phase), phenomenally 
increase the hardness and strength of the alloy, suppress coarsening of 
the fine crystal at high temperatures, and impart a resistance to heat. 
Furthermore, an atomic percentage for element M of less than 5% is 
undesirable, as this reduces the strength and hardness of the alloy. On 
the other hand, an atomic percentage exceeding 35% is also undesirable as 
this results in intermetallic compounds forming easily, which in turn lead 
to embrittlement of the alloy. 
Element R is one or more elements selected from the group consisting of Y, 
Ce, La, Nd and Mm (misch metal). 
In general, a misch metal mainly comprises La and/or Ce, and may also 
include additional complexes incorporating other rare earth metals, 
excluding the aforementioned La and Ce, as well as unavoidable impurities 
(Si, Fe, Mg, etc.). 
In particular, element R enhances the amorphous forming properties, and 
also raises the crystallization temperature of the amorphous phase. In 
this manner, the anti-corrosive properties can be improved, and the 
amorphous phase can be stabilized up to a high temperature. In addition, 
under the fine crystalline alloy manufacturing conditions, element R 
coexists with element M, and stabilizes the fine crystalline phase. 
Furthermore, an atomic percentage of element R exceeding 0.4% is 
undesirable as this results in the alloy being easily oxidized in addition 
to increased costs. 
In the aforementioned aluminium-based alloy according to the present 
invention represented by the formula Al.sub.x NiyM'.sub.z, the atomic 
percentages of Al, Ni, and element M' are restricted to 50-95%, 0.5-35% 
and 0.5-20%, respectively. This is due to the fact that when the 
composition of any of the aforementioned elements fall outside these 
specified ranges, it becomes difficult to form an amorphous component, as 
well as a supersaturated solid solution in which the amount of solute 
exceeds the critical solid solubility; this, in turn, results in the 
objective of the present invention, an aluminum-based alloy having an 
amorphous structure, an amorphous/fine crystalline complex structure or a 
fine crystalline structure, being unobtainable using an industrial 
quenching process incorporating a liquid quenching method. 
In addition, when diverging from the aforementioned composition ranges, it 
becomes difficult to obtain an amorphous phase for use in producing the 
fine crystalline complex structure, through crystallization of the 
amorphous phase produced by the quenching method using an appropriate 
heating process, or temperature control of a powder molding process which 
utilizes conventional powder metallurgy technology. 
An atomic percentage for Al of less than 50% is undesirable, as this 
results in significant embrittlement of the alloy. On the other hand, an 
atomic percentage for Al exceeding 95% is also undesirable, as this 
results in reduction of the strength and hardness of the alloy. 
Additionally, in the aforementioned composition ratio, the atomic 
percentage for Ni is within the range of 0.5-35%. If the incorporated 
amount of Ni is less than 0.5%, the strength and hardness of the alloy are 
reduced. On the other hand, an atomic percentage exceeding 35% results in 
intermetallic compounds forming easily, which in turn leads to 
embrittlement of the alloy. Thus both of these situations are undesirable. 
Furthermore, in the aforementioned composition ratio, the atomic percentage 
for element M' lies within the range of 0.5-20%. As in the aforementioned, 
if the incorporated amount of M' is less than 0.5%, the strength and 
hardness of the alloy are reduced. While, on the other hand, an atomic 
percentage exceeding 20% results in embrittlement of the alloy. Both of 
these situations are likewise undesirable. 
Element M' coexists with other elements, and improves the amorphous forming 
properties, in addition to raising the crystallization temperature of the 
amorphous phase. Most importantly, this element phenomenally improves the 
hardness and strength of the amorphous phase. As well, under the fine 
crystal manufacturing conditions, element M' also stabilizes the fine 
crystalline phase, forms stable or metastable intermetallic compounds with 
aluminum or other additional elements, disperses uniformly in the aluminum 
matrix (.alpha.-phase), phenomenally increases the hardness and strength 
of the alloy, suppresses coarsening of the fine crystal at high 
temperatures, and imparts a resistance to heat. 
In the aforementioned aluminium-based alloys according to the present 
invention represented by the formulae Al.sub.x Co.sub.y M"z and Al.sub.a 
Fe.sub.b L.sub.c, by adding predetermined amounts of Co and/or Fe to Al, 
the effect of quenching is enhanced, the amorphous and fine crystalline 
phases are more easily obtained, and the thermal stability of the overall 
structure is improved. In addition, the strength and hardness of the 
resulting alloy are also increased. 
In addition, by adding predetermined amounts of Mn and/or Cu to alloys 
consisting essentially of Al--Co.sub.2 or Al--Fe.sub.2, the strength and 
hardness of these alloys may be further improved. 
Furthermore, by adding predetermined amounts of Ti and/or Zr, the effect of 
quenching is enhanced, the amorphous and fine crystalline phases are more 
easily obtained, and the thermal stability of the overall structure is 
improved. 
The atomic percentage of Al is in the 50-95% range. An atomic percentage 
for Al of less than 50% is undesirable, as this results in embrittlement 
of the alloy. On the other hand, an atomic percentage for Al exceeding 95% 
is also undesirable, as this results in reduction of the strength and 
hardness of the alloy. 
Correspondingly, the atomic percentage of Co and/or Fe lies in the 0.5-35% 
range. When the atomic percentage of the aforementioned falls below 0.5%, 
the strength and hardness are not improved, while, on the other hand, when 
this atomic percentage exceeds 35%, embrittlement is observed, and the 
strength and toughness are reduced. Furthermore, in the case when Fe is 
added to an alloy comprising Al--Co.sub.2, if the atomic percentage 
exceeds 20%, embrittlement of the alloy begins to occur. 
The atomic percentage of Mn (manganese) and/or Cu (copper) lies in the 
0.5-20% range. When the atomic percentage of the aforementioned falls 
below 0.5%, improvements in the strength and hardness are not observed, 
while, on the other hand, when this atomic percentage exceeds 20%, 
embrittlement occurs, and the strength and toughness are reduced. 
The atomic percentage of Ti (titanium) and/or Zr (zirconium) lies in the 
range of up to one-half the atomic percentage of element M" or L. When the 
aforementioned atomic percentage is less than 0.5%, the quench effect is 
not improved, and, in the case when a crystalline state is incorporated 
into the alloy composition, the crystalline grains are not finely 
crystallized. On the other hand, when this atomic percentage exceeds 10%, 
embrittlement occurs, and toughness is reduced. In addition, the melting 
point rises, and melting become difficult to achieve. Furthermore, the 
viscosity of the liquid-melt increases, and thus, at the time of 
manufacturing, it becomes difficult to discharge this liquid-melt from the 
nozzle. 
In addition, when Ti or Zr is substituted in an amount exceeding one-half 
of the specified amount of element M", the hardness, strength and 
toughness are accordingly reduced. 
All of the aforementioned aluminum-based alloys according to the present 
invention can be manufactured by quench solidification of the alloy 
liquid-melts having the aforementioned compositions using a liquid 
quenching method. 
This liquid quenching method essentially entails rapid cooling of the 
melted alloy. Single roll, double roll, and submerged rotational spin 
methods have proved to be particularly effective. In these aforementioned 
methods, a cooling rate of 10.sup.4 to 10.sup.6 K/sec is easily 
obtainable. 
In order to manufacture a thin tape (alloy) using the aforementioned single 
or double roll methods, the liquid-melt is first poured into a storage 
vessel such as a silica tube, and then discharged, via a nozzle aperture 
at the tip of the silica tube, towards a copper roll of diameter 30 to 300 
mm, which is rotating at a fixed velocity in the range of 300 to 1000 rpm. 
In this manner, various types of thin tapes of thickness 5-500 .mu.m and 
width 1-300 mm can be easily obtained. 
On the other hand, fine wire-thin material can be easily obtained through 
the submerged rotational spin method by discharging the liquid-melt in 
order to quench it, via the nozzle aperture, into a refrigerant solution 
layer of depth 1 to 10 cm, maintained by means of centrifugal force inside 
an air drum rotating at 50 to 500 rpm, under argon gas back pressure. In 
this case, the angle between the liquid-melt discharged from the nozzle, 
and the refrigerant surface is preferably 60.degree. C. to 90.degree. C., 
and the relative velocity ratio of the the liquid-melt and the refrigerant 
surface is preferably 0.7 to 0.9. 
In addition, thin layers of aluminum-based alloy of the aforementioned 
compositions can also be obtained without using the above methods, by 
employing layer formation processes such as the sputtering method. In 
addition, aluminum alloy powder of the aforementioned compositions can be 
obtained by quenching the liquid-melt using various atomizer and spray 
methods such as a high pressure gas spray method. In the following, 
examples of structural states of the aluminum alloy obtained using the 
aforementioned methods are listed. 
(1) Non-crystalline phase; 
(2) Multiphase structure comprising an amorphous/Al fine crystalline phase; 
(3) Multiphase structure comprising an amorphous/stable or metastable 
intermetallic compound phase; 
(4) Multiphase structure comprising an Al/stable or metastable 
intermetallic compound or amorphous phase; and 
(5) Solid solution comprising a matrix of Al. 
The fine crystalline phase of the present invention represents a 
crystalline phase in which the crystal particles have an average maximum 
diameter of 1 .mu.m. The properties of the alloys possessing the 
aforementioned structural states are described in the following. 
An alloy of the structural state (amorphous phase) described in (1) above 
has a high strength, superior bending ductility, and a high toughness. 
Alloys possessing the structural phases (multiphase structures) described 
in (2) and (3) above have a high strength which is greater than that of 
the alloys of (amorphous) structural state (1) by a factor of 1.2 to 1.5. 
Alloys possessing the structural phases (multiphase structure and solid 
solution) described in (4) and (5) above have a greater toughness and 
higher strength than that of the alloys of structural states (1), (2) and 
(3). 
Each of the aforementioned structural states can be determined by a normal 
X-ray diffraction method or by observation using a transmission electron 
microscope. 
In the case of an amorphous phase, a halo pattern characteristic of this 
amorphous phase is evident. In the case of a multiphase structure 
comprising an amorphous/fine crystalline phase, a diffraction pattern 
formed from a halo pattern and characteristic diffraction peak, attributed 
to the fine crystalline phase, is displayed. In the case of a multiphase 
structure comprising an amorphous/intermetallic compound phase, a pattern 
formed from a halo pattern and characteristic diffraction peak, attributed 
to the intermetallic compound phase, is displayed. 
These amorphous and fine crystalline substances, as well as, amorphous/fine 
crystalline complexes can be obtained by means of various methods such as 
the aforementioned single and double roll methods, submerged rotational 
spin method, sputtering method, various atomizer methods, spray method, 
mechanical alloying method and the like. 
In addition, the amorphous/fine crystalline multiphase can be obtained by 
selecting the appropriate manufacturing conditions as necessary. 
By regulating the cooling rate of the alloy liquid-melt, any of the 
structural states described in (1) to (3) above can be obtained. 
By quenching the alloy liquid-melt of the Al-rich structure (e.g. 
structures with an Al atomic percentage of 92% or greater), any of the 
structural states described in (4) and (5) can be obtained. 
Subsequently, when the aforementioned amorphous phase structure is heated 
above a specific temperature, it decomposes to form crystal. This specific 
temperature is referred to as the crystallization temperature. 
By utilizing this heat decomposition of the amorphous phase, a complex of 
an aluminum solid solution phase in the fine crystalline state and 
different types of intermetallic compounds, determined by the alloy 
compositions therein, can be obtained. 
The aluminum-based alloy of the present invention displays superplasticity 
at temperatures near the crystallization temperature (crystallization 
temperature .+-.100.degree. C.), as well as, at the high temperatures 
within the fine crystalline stable temperature range, and thus processes 
such as extruding, pressing and hot forging can easily be performed. 
Consequently, aluminum-based alloys of the above-mentioned compositions 
obtained in the aforementioned thin tape, wire, plate and/or powder states 
can be easily formed into bulk materials by means of extruding, pressing 
and hot forging processes at the aforementioned temperatures. Furthermore, 
the aluminum-based alloys of the aforementioned compositions possess a 
high ductility, thus bending of 180.degree. is also possible. 
As well, the aluminum-based alloys having an amorphous phase or an 
amorphous/fine crystalline multiphase structure according to the present 
invention do not display structural or chemical non-uniformity of crystal 
grain boundary, segregation and the like, as seen in crystalline alloys. 
These alloys cause passivation due to formation of an aluminum oxide 
layer, and thus display a high resistance to corrosion. 
In particular, disadvantages exist when incorporating rare earth elements: 
due to the activity of these rare earth elements, non-uniformity occurs 
easily in the passive layer on the alloy surface resulting in the progress 
of corrosion from this portion towards the interior. However, since the 
alloys of the present invention do not incorporate rare earth elements, 
these aforementioned problems are effectively circumvented. 
In regards to the aluminum-based alloy of the present invention, the 
manufacturing of bulk-shaped (mass) material will now be explained. 
When heating the aluminum-based alloy according to the present invention, 
precipitation and crystallization of the fine crystalline phase is 
accompanied by precipitation of the aluminum matrix (.alpha.-phase), and 
when further heating beyond this temperature, the intermetallic compound 
also precipitates. Utilizing this property, bulk material possessing a 
high strength and ductility can be obtained. 
Concretely, the tape alloy manufactured by means of the aforementioned 
quench process is pulverized in a ball mill, and then powder pressed in a 
vacuum hot press under vacuum (e.g. 10.sup.-3 Torr) at a temperature 
slightly below the crystallization temperature (e.g. approximately 470K), 
thereby forming a billet for use in extruding with a diameter and length 
of several centimeters. This billet is set inside a container of an 
extruder, and is maintained at a temperature slightly greater than the 
crystallization temperature for several tens of minutes. Extruded 
materials can then be obtained in desired shapes such as round bars, etc. 
by extruding. 
Consequently, the aluminum-based alloy according to the present invention 
is useful as materials with a high strength, hardness and resistance to 
corrosion. Furthermore, it is possible to improve the mechanical 
properties by heat treatment; this alloy also stands up well to bending, 
and thus possesses superior properties such as the ability to be 
mechanically processed. 
In this manner, based on the aforementioned, the aluminum-based alloys 
according to the present invention can be used in a wide range of 
applications such as in aircraft, vehicles and ships, as well as, in the 
structural material for the engine portions thereof. In addition, the 
aluminum-based alloys of the present invention may also be employed as 
sash, roofing material and exterior material for use in construction, or 
as material for use in sea water equipment, nuclear reactors, and the like 
.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS 
First Preferred Embodiment! 
A molten alloy having a predetermined composition Al.sub.x M.sub.y R.sub.z 
was manufactured using a high frequency melting furnace. As shown in FIG. 
1, this melt was poured into a silica tube 1 with a small aperture 5 
(aperture diameter: 0.2 to 0.5 mm) at the tip, and then heat dissolved, 
following which the aforementioned silica tube 1 was positioned directly 
above copper roll 2. This roll 2 was then rotated at a high speed of 4000 
rpm, and argon gas pressure (0.7 kg/cm.sup.3) was applied to silica tube 
1. Quench solidification was subsequently performed by discharging the 
liquid-melt 3 from small aperture 5 of silica tube 1 onto the surface of 
roll 2 and quenching to yield an alloy tape 4. 
Under these manufacturing conditions, the numerous alloy tape samples 
(width: 1 mm, thickness: 20 .mu.m) of the compositions (atomic 
percentages) shown in Tables 1 and 2 were formed. Each sample was observed 
by both X-ray diffraction and TEM (transmission electron microscope). 
These results, shown in the structural state column of Tables 1 and 2, 
confirmed that an amorphous single-phase structure, a crystalline 
structure formed from an intermetallic compound or solid solution, and a 
two-phase structure (fcc-Al+Amo) formed by dispersing fine crystal grains, 
modified from aluminum having an fcc structure, into the amorphous matrix 
layer, were obtained. 
Subsequently, the hardness (Hv) and tensile rupture strength (.sigma.f: 
MPa) of each alloy tape sample were measured. These results are similarly 
shown in Tables 1 and 2. The hardness value (DPN: Diamond Pyramid Number) 
was measured according to the minute Vickers hardness scale. 
Additionally, a 180.degree. contact bending test was conducted by bending 
each sample 180.degree. and contacting the ends thereby forming a U-shape. 
The results of these tests are also shown in Tables 1 and 2: those samples 
which displayed ductility and did not rupture are designated Duc 
(ductile), while those which ruptured are designated Bri (brittle). 
It is clear from the results shown in Tables 1 and 2 that an aluminum-based 
alloy possessing a high bearing force and hardness, which endured bending 
and could undergo processing, was obtainable when the atomic percentages 
satisfied the relationships of 64.5.ltoreq.Al.ltoreq.95, 
5.ltoreq.M.ltoreq.35, and 0&lt;R.ltoreq.0.4. 
In contrast to normal aluminum-based alloys which possess an Hv of 
approximately 50 to 100 DPN, the samples according to the present 
invention, shown in Tables 1 and 2, display an extremely high hardness 
from 260 to 340 DPN. 
In addition, in regards to the tensile rupture strength (.sigma.f), normal 
age hardened type aluminum-based alloys (Al--Si--Fe type) possess values 
from 200 to 600 MPa, however, the samples according to the present 
invention have clearly superior values in the range from 800 to 1250 MPa. 
Furthermore, when considering that the tensile strengths of aluminum-based 
alloys of the AA6000 series (alloy name according to the Aluminum 
Association (U.S.A.)) and AA7000 series which lie in the range from 250 to 
300 MPa, Fe-type structural steel sheets which possess a value of 
approximately 400 MPa, and high tensile strength steel sheets of Fe-type 
which range from 800 to 980 MPa, it is clear that the aluminum-based 
alloys according to the present invention display superior values. 
FIG. 2 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.88 Ni.sub.11.6 Ce.sub.0.4. In this FIG., 
the crystal peak (not discernible) appears as a broad peak pattern with 
the alloy sample displaying an amorphous single phase structure. 
FIG. 3 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.89.7 Ni.sub.5 Fe.sub.5 Ce.sub.0.3. In 
this FIG., a two-phase structure is displayed in which fine Al particles 
having an fcc structure of the nano-scale are dispersed into the amorphous 
phase. In the FIG., (111) and (200) display the crystal peaks of Al having 
an fcc structure. 
FIG. 4 shows the DSC (Differential Scanning Calorimetry) curve in the case 
when an alloy having the composition of Al.sub.89.6 Ni.sub.5 Co.sub.5 
Ce.sub.0.4 is heated at an increase temperature rate of 0.67 K/s. 
FIG. 5 shows the DSC curve in the case when an alloy having the composition 
of Al.sub.88 Ni.sub.11.6 Y.sub.0.4 is heated at an increase temperature 
rate of .sub.0.67 K/s. 
As is clear from FIGS. 4 and 5, the broad peak appearing at lower 
temperatures represents the crystallization peak of Al particles having an 
fcc structure, while the sharp peak at higher temperatures represents the 
crystallization peak of the alloys. Due to the existence of these two 
peaks, when performing heat treatment such as quench hardening at an 
appropriate temperature, the volume percentage of the Al particles 
dispersed into the amorphous matrix phase can be controlled. As a result, 
it is clear that the mechanical properties can be improved through heat 
treatment. 
In addition, in order to show criticality of the aforementioned composition 
ratios of 64.5.ltoreq.Al.ltoreq.95, 5.ltoreq.M .ltoreq.35, and 
0&lt;R.ltoreq.0.4, FIGS. 6-14 are provided. 
The graph in FIG. 6 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x M.sub.99.7-x 
Y.sub.0.3 (in which element M is Ti, V, Cr, or Mn) corresponding to 
various values of x. 
The graph in FIG. 7 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x M.sub.99.7-x 
Ce.sub.0.3 (in which element M is Fe, Ni, Co, or Cu) corresponding to 
various values of x. 
The graph in FIG. 8 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x M.sub.99.7-x 
La.sub.0.3 (in which element M is Zr, Nb, or Mo) corresponding to various 
values of x. 
According to the graphs of FIGS. 6-8, it can be seen that an alloy having a 
composition of Al.sub.x M.sub.y R.sub.z in which the atomic percentage for 
Al is less than 64.5% or exceeds 95% is undesirable, since such an alloy 
may not have sufficient strength. 
The graph in FIG. 9 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.99.6-y M.sub.y 
Ce.sub.0.4 (in which element M is Ti, V, Cr, or Mn) corresponding to 
various values of y. 
The graph in FIG. 10 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.99.6-y M.sub.y 
Nd.sub.0.4 (in which element M is Fe, Ni, Co, or Cu) corresponding to 
various values of y. 
The graph in FIG. 11 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.99.6-y M.sub.y 
Mm.sub.0.4 (in which element M is Zr, Nb, or Mo) corresponding to various 
values of y. 
According to the graphs of FIGS. 9-11, it can be seen that an alloy having 
a composition of Al.sub.x M.sub.y R.sub.z in which the atomic percentage 
for element M is less than 5% or exceeds 35% is undesirable, since such an 
alloy may not have sufficient strength. 
The graph in FIG. 12 shows variation of the corrosion rate (in 1N-HCl 
solution) of alloys having the compositions of Al.sub.89-z M.sub.11 
Y.sub.z (in which element M is Ti, V, Cr, or Mn) corresponding to various 
values of z. 
The graph in FIG. 13 shows variation of the corrosion rate (in 1N-HCl 
solution) of alloys having the compositions of Al.sub.89-z M.sub.11 
Nd.sub.z (in which element M is Fe, Ni, Co, or Cu) corresponding to 
various values of z. 
The graph in FIG. 14 shows variation of the corrosion rate (in 1N-HCl 
solution) of alloys having the compositions of Al.sub.89-z M.sub.11 
La.sub.z (in which element M is Zr, Nb, or corresponding to various values 
of z. 
According to the graphs of FIGS. 12-14, it can be seen that an alloy having 
a composition of Al.sub.x M.sub.y R.sub.z in which the atomic percentage 
for element R exceeds 0.4% is undesirable, since such an alloy may corrode 
easily. 
Second Preferred Embodiment! 
In a manner similar to the first preferred embodiment, a molten alloy 
having a predetermined composition Al.sub.x Ni.sub.y M'.sub.z was 
manufactured using a high frequency melting furnace. As shown in FIG. 1, 
this melt was poured into a silica tube 1 with a small aperture 5 
(aperture diameter: 0.2 to 0.5 mm) at the tip, and then heat dissolved, 
following which the aforementioned silica tube 1 was positioned directly 
above copper roll 2. This roll 2 was then rotated at a high speed of 4000 
rpm, and argon gas pressure (0.7kg/cm.sup.3) was applied to silica tube 1. 
Quench solidification was subsequently performed by discharging the 
liquid-melt from small aperture 5 of silica tube 1 onto the surface of 
roll 2 and quenching to yield an alloy tape 4. 
Under these manufacturing conditions, the numerous alloy tape samples 
(width: 1 mm, thickness: 20 .mu.m) of the compositions (atomic 
percentages) shown in Tables 3 and 4 were formed. Each sample was observed 
by both X-ray analysis and TEM (transmission electron microscope). 
These results, shown in the structural state column of Tables 3 and 4, 
confirmed that an amorphous single-phase structure, a crystalline 
structure formed from an intermetallic compound or solid solution, and a 
two-phase structure (fcc-Al+Amo) formed by dispersing fine crystal grains, 
modified from aluminum having an fcc structure, into the amorphous matrix 
layer, were obtained. 
Subsequently, the hardness (Hv) and tensile rupture strength (.sigma.f: 
MPa) of each alloy tape sample were measured. These results are similarly 
shown in Tables 3 and 4. The hardness value (DPN: Diamond Pyramid Number) 
was measured according to the minute Vickers hardness scale. 
Additionally, the 180.degree. contact bending test was conducted by bending 
each alloy tape sample 180.degree. and contacting the ends thereby forming 
a U-shape. 
The results of these tests are also shown in Tables 3 and 4: those samples 
which displayed ductility and did not rupture are designated Duc 
(ductile), while those which ruptured are designated Bri (brittle). 
It is clear from the results shown in Tables 3 and 4 that an aluminum-based 
alloy possessing a high bearing force and hardness, which endured bending 
and could undergo processing, was obtainable when the atomic percentages 
satisfied the relationships of 50.ltoreq.Al.ltoreq.95, 
0.5.ltoreq.Ni.ltoreq.35, and 0.5.ltoreq.M'.ltoreq.20. 
In contrast to normal aluminum-based alloys which possess an Hv of 
approximately 50 to 100 DPN, the samples according to the present 
invention shown in Tables 3 and 4 display an extremely high hardness 
ranging from 260 to 400 DPN. 
In addition, in regards to the tensile rupture strength (.sigma.f), normal 
age hardened type aluminum-based alloys (Al--Si--Fe type) possess values 
from 200 to 600 MPa, however, the samples according to the present 
invention have clearly superior values in the range from 780 to 1150 MPa. 
Furthermore, when considering that the tensile strengths of aluminum-based 
alloys of the AA6000 series and AA7000 series which lie in the range from 
250 to 300 MPa, Fe-type structural steel sheets which possess a value of 
approximately 400 MPa, and high tensile strength steel sheets of Fe-type 
which range from 800 to 980 MPa, it is clear that the aluminum-based 
alloys according to the present invention display superior values. 
FIG. 15 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.87 Ni.sub.12 Mn.sub.1. In this FIG., the 
crystal peak (not discernible) appears as a broad peak pattern with the 
alloy sample displaying an amorphous single phase structure. 
FIG. 16 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.88 Ni.sub.9 Co.sub.3. In this FIG., a 
two-phase structure is displayed in which fine Al particles having an fcc 
structure of the nano-scale are dispersed into the amorphous phase. In the 
FIG., (111) and (200) display the crystal peaks of Al having an fcc 
structure. 
FIG. 17 shows the DSC (Differential Scanning Calorimetry) curve in the case 
when an alloy having the composition of Al88Ni.sub.11 Zr.sub.1 is heated 
at an increase temperature rate of 0.67 K/s. 
FIG. 18 shows the DSC curve in the case when an alloy having the 
composition of Al.sub.88 Ni.sub.11 Fe.sub.1 is heated at an increase 
temperature rate of 0.67 K/s. 
As is clear from FIGS. 17 and 18, the broad peak appearing at lower 
temperatures represents the crystallization peak of Al particles having an 
fcc structure, while the sharp peak at higher temperatures represents the 
crystallization peak of the alloys. Due to the existence of these two 
peaks, when performing heat treatment such as quench hardening at an 
appropriate temperature, the volume percentage of the Al particles 
dispersed into the amorphous matrix phase can be controlled. As a result, 
it is clear that the mechanical properties can be improved through heat 
treatment. 
In addition, in order to show criticality of the aforementioned composition 
ratios of 50.ltoreq.Al.ltoreq.95, 0.5.ltoreq.Ni .ltoreq.35, and 
0.5.ltoreq.M'.ltoreq.20, FIGS. 19-24 are provided. 
The graph in FIG. 19 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x Ni.sub.96-x 
M'.sub.4 and Al.sub.x Ni.sub.85-x M'.sub.15 (in which element M' is Ti, V, 
Cr, or Mn) corresponding to various values of x. 
The graph in FIG. 20 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x Ni.sub.96-x 
M'.sub.4 and Al.sub.x Ni.sub.85-x M'.sub.15 (in which element M' is Co, 
Cu, or Zr) corresponding to various values of x. 
According to the graphs of FIGS. 19 and 20, it can be seen that an alloy 
having a composition of Al.sub.x Ni.sub.y M'.sub.z in which the atomic 
percentage for Al is less than 50% or exceeds 95% is undesirable, since 
such an alloy may not have sufficient strength. 
The graph in FIG. 21 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-y Ni.sub.y 
M'.sub.15 (in which element M' is Ti, V, Mn, or Fe) corresponding to 
various values of y. 
The graph in FIG. 22 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-y Ni.sub.y 
M'.sub.15 (in which element M' is Co, Cu, or Zr) corresponding to various 
values of y. 
According to the graphs of FIGS. 21 and 22, it can be seen that an alloy 
having a composition of Al.sub.x Ni.sub.y M'.sub.z in which the atomic 
percentage for Ni is less than 0.5% or exceeds 35% is undesirable, since 
such an alloy may not have sufficient strength. 
The graph in FIG. 23 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-z Ni.sub.15 
M'.sub.z (in which element M' is Ti, V, Mn, or Fe) corresponding to 
various values of z. 
The graph in FIG. 24 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-z Ni.sub.15 
M'.sub.z (in which element M' is Co, Cu, or Zr) corresponding to various 
values of z. 
According to the graphs of FIGS. 23 and 24, it can be seen that an alloy 
having a composition of Al.sub.x Ni.sub.y M'.sub.z in which the atomic 
percentage for element M' is less than 0.5% or exceeds 20% is undesirable, 
since such an alloy may not have sufficient strength. 
Third Preferred Embodiment! 
In a manner similar to the first and second preferred embodiments, a molten 
alloy having a predetermined composition Al.sub.x Co.sub.y M".sub.z or 
Al.sub.a Fe.sub.b L.sub.c was manufactured using a high frequency melting 
furnace. As shown in FIG. 1, this melt was poured into a silica tube 1 
with a small aperture 5 (aperture diameter: 0.2 to 0.5 mm) at the tip, and 
then heat dissolved, following which the aforementioned silica tube 1 was 
positioned directly above copper roll 2. This roll 2 was then rotated at a 
high speed of 4000 rpm, and argon gas pressure (0.7kg/cm.sup.3) was 
applied to silica tube 1. Quench solidification was subsequently performed 
by discharging the liquid-melt from small aperture 5 of silica tube 1 onto 
the surface of roll 2 and quenching to yield an alloy tape 4. 
Under these manufacturing conditions, the numerous alloy tape samples 
(width: 1 mm, thickness: 20 .mu.m) of the compositions (atomic 
percentages) shown in Tables 5 to 7 were formed. Each sample was observed 
by both X-ray diffraction and TEM (transmission electron microscope). 
These results, shown in the structural state column of Tables 5 to 7, 
confirmed that an amorphous (Amo) single-phase structure, a crystalline 
structure (Com) formed from an intermetallic compound or solid solution, a 
multiphase structure (fcc-Al+Amo) formed from fine crystal grains of 
aluminum having an fcc structure, and a structure formed from the 
aforementioned amorphous and crystalline structures, were obtained. 
Subsequently, the hardness (Hv) and tensile rupture strength (.sigma.f: 
MPa) of each alloy tape sample were measured. These results are similarly 
shown in Tables 5 to 7. The hardness value (DPN: Diamond Pyramid Number) 
was measured according to the minute Vickers hardness scale. 
Additionally, the 180.degree. contact bending test was conducted by bending 
each sample 180.degree. and contacting the ends thereby forming a U-shape. 
The results of these tests are also shown in Tables 5 to 7: those samples 
which displayed ductility and did not rupture are designated Duc 
(ductile), while those which did rupture are designated Bri (brittle) . 
It is clear from the results shown in Tables 5 to 7 that when element M" is 
added to a Al--Co.sub.2 --component alloy, wherein M" is one or more 
elements selected from the group consisting of Mn, Fe and Cu, an 
aluminum-based alloy possessing a high bearing force and hardness, which 
endured bending and could undergo processing, was obtainable when the 
atomic percentages satisfied the relationships of 50.ltoreq.Al.ltoreq.95, 
0.5.ltoreq.Co.ltoreq.35, and 0.5.ltoreq.M".ltoreq.20. 
Furthermore it is also clear from the results shown in Tables 5 to 7 that 
when element L is added to a Al--Fe.sub.2 --component alloy, wherein L is 
one or more elements selected from the group consisting of Mn and Cu, an 
aluminum-based alloy possessing a high bearing force and hardness, which 
endured bending and could undergo processing, was obtainable when the 
atomic percentages satisfied the relationships of 50.ltoreq.Al.ltoreq.95, 
0.5.ltoreq.Fe.ltoreq.35, and 0.5.ltoreq.L.ltoreq.20. 
In contrast to normal aluminum-based alloys which possess an Hv of 
approximately 50 to 100 DPN, the samples according to the present 
invention shown in Tables 5 and 7 display an extremely high hardness 
ranging from 165 to 387 DPN. 
In addition, in regards to the tensile rupture strength (.sigma.f), normal 
age hardened type aluminum-based alloys (Al--Si--Fe type) possess values 
from 200 to 600 MPa, however, the samples according to the present 
invention have clearly superior values in the range from 760 to 1270 MPa. 
Furthermore, when considering that the tensile strengths of aluminum-based 
alloys of the AA6000 series and AA7000 series which lie in the range from 
250 to 300 MPa, Fe-type structural steel sheets which possess a value of 
approximately 400 MPa, and high tensile strength steel sheets of Fe-type 
which range from 800 to 980 MPa, it is clear that the aluminum-based 
alloys according to the present invention display superior values. 
FIG. 25 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.89 Co.sub.8 Mn.sub.3. In this FIG., the 
crystal peak (not discernible) appears as a broad peak pattern with the 
alloy sample displaying an amorphous single phase structure. 
FIG. 26 shows the analysis result of the X-ray diffraction of an alloy 
having the composition of Al.sub.90 Co.sub.6 Fe.sub.4. In this FIG., a 
multiphase structure is displayed which comprises an amorphous phase and a 
fine Al crystalline phase having an fcc structure of the nanoscale. In the 
FIG., (111) and (200) display the crystal peaks of Al having an fcc 
structure. 
FIG. 27 shows the DSC (Differential Scanning Calorimetry) curve in the case 
when an alloy having the composition of Al.sub.90 Co.sub.9 Cu.sub.1 is 
heated at an increase temperature rate of 0.67 K/s. 
FIG. 28 shows the DSC curve in the case when an alloy having the 
composition of Al.sub.90 Co.sub.9 Mn.sub.1 is heated at an increase 
temperature rate of 0.67 K/s. 
As is clear from FIGS. 27 and 28, the broad peak appearing at lower 
temperatures represents the crystallization peak of Al particles having an 
fcc structure, while the sharp peak at higher temperatures represents the 
crystallization peak of the alloys. Due to the existence of these two 
peaks, when performing heat treatment such as quench hardening at an 
appropriate temperature, the volume percentage of the Al particles 
dispersed into the amorphous matrix phase can be controlled. As a result, 
it is clear that the mechanical properties can be improved through heat 
treatment. 
In addition, in order to show criticality of the aforementioned composition 
ratios of 50.ltoreq.Al.ltoreq.95, 0.5.ltoreq.Co .ltoreq.35, and 
0.5.ltoreq.M".ltoreq.20 for Al.sub.x Co.sub.y M".sub.z, or of 
50.ltoreq.Al.ltoreq.95, 0.5.ltoreq.Fe.ltoreq.35, and 
0.5.ltoreq.L.ltoreq.20 for Al.sub.a Fe.sub.b L.sub.c, FIGS. 29-38 are 
provided. 
The graph in FIG. 29 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.x Co.sub.96-x 
M".sub.4 and Al.sub.x Co.sub.85-x M".sub.15 (in which element M" is Mn, 
Fe, or Cu) corresponding to various values of x. According to this graph, 
it can be seen that an alloy having a composition of Al.sub.x Co.sub.y 
M".sub.z in which the atomic percentage for Al is less than 50% or exceeds 
95% is undesirable, since such an alloy may not have sufficient strength. 
The graph in FIG. 30 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-y Co.sub.y 
M".sub.15 (in which element M" is Mn, Fe, or Cu) corresponding to various 
values of y. According to this graph, it can be seen that an alloy having 
a composition of Al.sub.x Co.sub.y M".sub.z in which the atomic percentage 
for Co is less than 0.5% or exceeds 35% is undesirable, since such an 
alloy may not have sufficient strength. 
The graph in FIG. 31 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-z Co.sub.15 
M".sub.z (in which element M" is Mn, Fe, or Cu) corresponding to various 
values of z. According to this graph, it can be seen that an alloy having 
a composition of Al.sub.x Co.sub.y M".sub.z in which the atomic percentage 
for element M" is less than 0.5% or exceeds 20% is undesirable, since such 
an alloy may not have sufficient strength. 
The graph in FIG. 32 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.a Fe.sub.97-a 
L.sub.3 and Al.sub.a Fe.sub.85-a L.sub.3 (in which L is Mn or Cu) 
corresponding to various values of a. According to this graph, it can be 
seen that an alloy having a composition of Al.sub.a Fe.sub.b L.sub.c in 
which the atomic percentage for Al is less than 50% or exceeds 95% is 
undesirable, since such an alloy may not have sufficient strength. 
The graph in FIG. 33 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-b Fe.sub.b 
L.sub.15 (in which L is Mn or Cu) corresponding to various values of b. 
According to this graph, it can be seen that an alloy having a composition 
of Al.sub.a Fe.sub.b L.sub.c in which the atomic percentage for Fe is less 
than 0.5% or exceeds 35% is undesirable, since such an alloy may not have 
sufficient strength. 
The graph in FIG. 34 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.85-c Fe.sub.15 
L.sub.c (in which L is Mn or Cu) corresponding to various values of c. 
According to this graph, it can be seen that an alloy having a composition 
of Al.sub.a Fe.sub.b L.sub.c in which the atomic percentage for L is less 
than 0.5% or exceeds 20% is undesirable, since such an alloy may not have 
sufficient strength. 
The graph in FIG. 35 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.88 Co.sub.6 
M".sub.6(1-a) Zr.sub.6a (in which element M" is Mn, Fe, or Cu) 
corresponding to various values of a. According to this graph, it can be 
seen that an alloy having a composition of Al.sub.x Co.sub.y M".sub.z in 
which a part of element M" is substituted by Zr but in which the atomic 
percentage for Zr exceeds onehalf of that of element M" is undesirable, 
since such an alloy may not have sufficient strength. 
The graph in FIG. 36 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.88 Co.sub.6 
M".sub.6(1-a) Ti.sub.6a (in which element M" is Mn, Fe, or Cu) 
corresponding to various values of a. According to this graph, it can be 
seen that an alloy having a composition of Al.sub.x Co.sub.y M".sub.z in 
which a part of element M" is substituted by Ti but in which the atomic 
percentage for Ti exceeds one-half of that of element M" is undesirable, 
since such an alloy may not have sufficient strength. 
The graph in FIG. 37 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.86 Fe.sub.8 
L.sub.6(1-x) Zr.sub.6x (in which L is Mn or Cu) corresponding to various 
values of x. According to this graph, it can be seen that an alloy having 
a composition of Al.sub.a Fe.sub.b L.sub.c in which a part of L is 
substituted by Zr but in which the atomic percentage for Zr exceeds 
one-half of that of L is undesirable, since such an alloy may not have 
sufficient strength. 
The graph in FIG. 38 shows variation of the tensile rupture strength 
(.sigma.f) of alloys having the compositions of Al.sub.86 Fe.sub.8 
L.sub.6(1-x) Ti.sub.6x (in which L is Mn or Cu) corresponding to various 
values of x. According to this graph, it can be seen that an alloy having 
a composition of Al.sub.a Fe.sub.b L.sub.c in which a part of L is 
substituted by Ti but in which the atomic percentage for Ti exceeds 
one-half of that of L is undesirable, since such an alloy may not have 
sufficient strength. 
Comparative Tests! 
U.S. Pat. No. 4,595,429 (Le Caer, et al.) discloses alloys having the 
composition Al.sub.a M.sub.b M'.sub.c X.sub.d Y.sub.e, in which: 
50.ltoreq.a.ltoreq.95 atom %; M representing one or more metals of the 
group Mn, Ni, Cu, Zr, Ti, C, Cr, Fe, and Co, with 0.ltoreq.b.ltoreq.40 
atom %; M' representing Mo and/or W, with 0.ltoreq.c.ltoreq.15 atom %; X 
representing one or more elements of the group Ca, Li, Mg, Ge, Si, and Zn, 
with 0.ltoreq.d.ltoreq.20 atom %; and Y representing impurities such as O, 
N, C, H, He, Ga, etc., the proportions of which does not exceed 3 atom %. 
In order to determine differences in bending ductility and tensile strength 
between the compositions of Le Caer, et al., and those of the present 
invention, several alloys were prepared and tested in accordance with 
anomalous X-ray scattering (AXS). The results are shown in FIGS. 39-42. 
Although the alloys according to Le Caer, et al., are similar in 
composition to the alloys according to the present invention, the alloys 
of Le Caer, et al., do not possess sufficient bending ductility or tensile 
strength. 
The graph in FIG. 39 shows structure-analysis data of an alloy according to 
Le Caer, et al., having the composition of Al.sub.70 Ge.sub.20 Ni.sub.10. 
It is noted that this composition corresponds to a composition of Le Caer, 
et al., in which a=70, M is Ni, b=10, c=0, X is Ge, d=20, and e=0. 
The graph in FIG. 40 shows structure-analysis data of an alloy according to 
Le Caer, et al., having the composition of Al.sub.70 Si.sub.15 Ni.sub.15. 
It is noted that this composition corresponds to a composition of Le Caer, 
et al., in which a=70, M is Ni, b=15, c=0, X is Si, d=15, and e=0. 
FIG. 41 is a graph showing structure-analysis data of an alloy according to 
the present invention having the composition of Al.sub.88.7 Ni.sub.11 
Ce.sub.0.3. 
FIG. 42 is a graph showing structure-analysis data of an alloy according to 
the present invention having the composition of Al.sub.88 Ni.sub.11 
Fe.sub.1. 
In these graphs in FIGS. 39-42, one axis (Q) represents the wave number 
vector, and the other axis (.DELTA.I(Q)) represents the differential 
intensity profile at incident energy. 
According to the graphs of FIGS. 39 and 40, the differential intensity 
profile values are partially negative, and this indicates the existence of 
a short periodical regular array of elements which produces a brittle 
amorphous structure. Accordingly, these alloys do not have bending 
ductility. In contrast, it can be seen from the graphs of FIGS. 41 and 42 
showing data of alloys according to the present invention that the 
differential intensity profile is always positive for any value of the 
wave number vector. This indicates that the amorphous structures of the 
alloys according to the present invention are homogeneous, on the whole, 
and the alloys exhibit bending ductility. This makes a test of the tensile 
strength possible and it is found that the alloys of the present invention 
possess a high strength of over 750 MPa and a desirable Vickers hardness 
in the range of 150-385. 
Although the invention has been described in detail herein with reference 
to its preferred embodiments and certain described alternatives, it is to 
be understood that this description is by way of example only, and it is 
not to be construed in a limiting sense. It is further understood that 
numerous changes in the details of the embodiments of the invention, and 
additional embodiments of the invention, will be apparent to, and may be 
made by persons of ordinary skill in the art having reference to this 
description. It is contemplated that all such changes and additional 
embodiments are within the spirit and true scope of the invention as 
claimed below. 
TABLE 1 
__________________________________________________________________________ 
Sample 
Alloy composition Structural 
Bending 
No. (at %) .sigma.f (MPa) 
Hv (DPN) 
state test 
__________________________________________________________________________ 
1 Al.sub.89.6 Ni.sub.5 Co.sub.5 Ce.sub.0.4 
1240 317 fcc-Al + Amo 
Duc 
2 Al.sub.88.7 Ni.sub.11 Nd.sub.0.3 
1170 305 fcc-Al + Amo 
Duc 
3 Al.sub.88.7 Ni.sub.11 La.sub.0.3 
1050 260 amorphous 
Duc 
4 Al.sub.88.7 Ni.sub.11 Ce.sub.0.3 
1030 272 amorphous 
Duc 
5 Al.sub.88.7 Cu.sub.11 Y.sub.0.3 
1190 310 fcc-Al + Amo 
Duc 
6 Al.sub.88.7 Mn.sub.11 Ce.sub.0.3 
910 307 fcc-Al + Amo 
Duc 
7 Al.sub.88.7 Fe.sub.11 Mn.sub.0.3 
900 298 fcc-Al + Amo 
Duc 
8 Al.sub.87.6 Ni.sub.11 Cr.sub.1 Y.sub.0.4 
800 340 fcc-Al + Amo 
Duc 
9 Al.sub.87.6 Ni.sub.11 V.sub.1 Y.sub.0.4 
840 305 amorphous 
Duc 
10 Al.sub.87.6 Ni.sub.11 Ti.sub.1 Y.sub.0.4 
1030 332 amorphous 
Duc 
11 Al.sub.87.6 Ni.sub.11 Zr.sub.1 Ce.sub.0.4 
960 280 amorphous 
Duc 
12 Al.sub.87.6 Ni.sub.11 Nb.sub.1 Ce.sub.0.4 
980 317 fcc-Al + Amo 
Duc 
13 Al.sub.87.6 Ni.sub.11 Mo.sub.1 Ce.sub.0.4 
1020 320 fcc-Al + Amo 
Duc 
__________________________________________________________________________ 
TABLE 2 
______________________________________ 
Sam- 
ple Alloy composition 
.sigma.f 
Hv Structural 
Bending 
No. (at %) (MPa) (DPN) state test 
______________________________________ 
14 Al.sub.60.7 Fe.sub.39 Y.sub.0.3 
--*.sup.1 
520 Crystalline 
Bri 
15 Al.sub.98.7 Fe.sub.1 Ce.sub.0.3 
440 120 fcc-Al Duc 
16 Al.sub.99.7 Ce.sub.0.3 
400 107 fcc-Al Duc 
17 Al.sub.60 Fe.sub.40 
--*.sup.1 
520 Crystalline 
Bri 
______________________________________ 
*.sup.1 Tensile test could not be conducted due to brittle nature. 
TABLE 3 
______________________________________ 
Sam- Alloy 
ple composition 
.sigma.f 
Hv Structural 
Bending 
No. (at %) (MPa) (DPN) state test 
______________________________________ 
18 Al.sub.88 Ni.sub.7 Co.sub.5 
1065 316 amorphous Duc 
19 Al.sub.88 Ni.sub.8 Co.sub.4 
1061 313 amorphous Duc 
20 Al.sub.88 Ni.sub.9 Co.sub.3 
996 307 amorphous Duc 
21 Al.sub.88 Ni.sub.10 Co.sub.2 
813 306 fcc-Al + Amo 
Duc 
22 Al.sub.88 Ni.sub.11 Co.sub.1 
931 295 fcc-Al + Amo 
Duc 
23 Al.sub.88 Ni.sub.8 Fe.sub.4 
1080 302 fcc-Al + Amo 
Duc 
24 Al.sub.88 Ni.sub.9 Fe.sub.3 
960 309 fcc-Al + Amo 
Duc 
25 Al.sub.88 Ni.sub.10 Fe.sub.2 
915 304 fcc-Al + Amo 
Duc 
26 Al.sub.88 Ni.sub.11 Fe.sub.1 
928 311 fcc-Al + Amo 
Duc 
27 Al.sub.88 Ni.sub.11 Cu.sub.1 
780 327 fcc-Al + Amo 
Duc 
28 Al.sub.88 Ni.sub.11 Mn.sub.1 
930 302 fcc-Al + Amo 
Duc 
29 Al.sub.88 Ni.sub.11 V.sub.1 
797 363 fcc-Al + Amo 
Duc 
30 Al.sub.88 Ni.sub.11 Ti.sub.1 
1047 368 fcc-Al + Amo 
Duc 
31 Al.sub.88 Ni.sub.11 Zr.sub.1 
954 276 fcc-Al + Amo 
Duc 
______________________________________ 
TABLE 4 
______________________________________ 
Alloy 
Sample composition 
.sigma.f 
Hv Structural 
Bending 
No. (at %) (MPa) (DPN) state test 
______________________________________ 
32 Al.sub.90 Ni.sub.5 Co.sub.5 
1150 380 fcc-Al + 
Duc 
Amo 
33 Al.sub.87 Ni.sub.12 Mn.sub.1 
953 262 amorphous 
Duc 
34 Al.sub.88 Ni.sub.7 V.sub.5 
1070 331 fcc-Al + 
Duc 
Amo 
35 Al.sub.95 Ni.sub.0.3 Cu.sub.4.7 
420 117 fcc-Al Duc 
36 Al.sub.95 Ni.sub.0.3 Cu.sub.4.7 
400 109 fcc-Al Duc 
37 Al.sub.95 Ni.sub.0.3 Fe.sub.4.7 
450 123 fcc-Al Duc 
38 Al.sub.88 Mn.sub.12 
--*.sup.1 
550 Crystalline 
Bri 
39 Al.sub.73 Ni.sub.2 Fe.sub.25 
--*.sup.1 
570 Crystalline 
Bri 
40 Al.sub.50 Ni.sub.40 Fe.sub.10 
--*.sup.1 
530 Crystalline 
Bri 
41 Al.sub.94.6 Ni.sub.5 Cu.sub.0.4 
380 102 fcc-Al Duc 
42 Al.sub.94 Ni.sub.6 
540 180 fcc-Al Duc 
43 Al.sub.96 Ni.sub.2 Co.sub.2 
400 120 fcc-Al Duc 
44 Al.sub.55 Ni.sub.40 Fe.sub.5 
--*.sup.1 
520 Crystalline 
Bri 
______________________________________ 
*.sup.1 Tensile test could not be conducted due to brittle nature. 
TABLE 5 
__________________________________________________________________________ 
Alloy 
composition 
(Subscript numerals 
Sample 
represent atomic 
.sigma.f 
Hv Structural 
Bending 
No. percentage) 
(MPa) 
(DPN) 
state test 
__________________________________________________________________________ 
45 Al.sub.98 Co.sub.1 Mn.sub.1 
400 
110 fcc-Al Duc Comparative example 
46 Al.sub.95 Co.sub.4 Mn.sub.1 
780 
215 fcc-Al Duc Example 
47 Al.sub.90 Co.sub.8 Mn.sub.2 
1270 
330 fcc-Al + Amo 
Duc Example 
48 Al.sub.80 Co.sub.15 Mn.sub.5 
1115 
315 fcc-Al + Amo 
Duc Example 
49 Al.sub.70 Co.sub.25 Mn.sub.5 
1210 
320 fcc-Al + Amo 
Duc Example 
50 Al.sub.60 Co.sub.30 Mn.sub.10 
980 
370 Amo + Com 
Duc Example 
51 Al.sub.50 Co.sub.30 Mn.sub.20 
960 
360 Amo + Com 
Duc Example 
52 Al.sub.45 Co.sub.35 Mn.sub.20 
-- 550 Com Bri Comparative example 
53 Al.sub.50 Co.sub.40 Mn.sub.10 
-- 490 Com Bri Comparative example 
54 Al.sub.60 Co.sub.35 Mn.sub.5 
960 
370 Amo + Com 
Duc Example 
55 Al.sub.65 Co.sub.30 Mn.sub.5 
975 
340 fcc-Al + Amo 
Duc Example 
56 Al.sub.70 Co.sub.20 Mn.sub.10 
1010 
340 fcc-Al + Amo 
Duc Example 
57 Al.sub.80 Co.sub.10 Mn.sub.10 
1015 
345 fcc-Al + Amo 
Duc Example 
58 Al.sub.96 Co.sub.1 Mn.sub.3 
760 
180 fcc-Al Duc Example 
59 Al.sub.95 Co.sub.0.5 Mn.sub.4.5 
760 
165 fcc-Al Duc Example 
60 Al.sub.94 Co.sub.0.3 Mn.sub.5.7 
445 
85 fcc-Al Duc Comparative example 
__________________________________________________________________________ 
TABLE 6 
__________________________________________________________________________ 
Alloy 
composition 
(Subscript numerals 
Sample 
represent atomic 
.sigma.f 
Hv Structural 
Bending 
No. percentage) 
(MPa) 
(DPN) 
state test 
__________________________________________________________________________ 
61 Al.sub.70 Co.sub.5 Mn.sub.25 
-- 520 Com Bri Comparative example 
62 Al.sub.72 Co.sub.8 Mn.sub.20 
1195 
360 Amo + Com 
Duc Example 
63 Al.sub..sub.80 Co.sub.10 Mn.sub.10 
1145 
320 fcc-Al + Amo 
Duc Example 
64 Al.sub.89 Co.sub.10 Mn.sub.1 
1230 
387 fcc-Al + Amo 
Duc Example 
65 Al.sub.91 Co.sub.8.5 Mn.sub.0.5 
1200 
330 fcc-Al + Amo 
Duc Example 
66 Al.sub.89 Co.sub.10.7 Mn.sub.0.3 
460 
120 fcc-Al + Amo 
Duc Comparative example 
67 Al.sub.98 Co.sub.1 Fe.sub.1 
420 
125 fcc-Al Duc Comparative example 
68 Al.sub.80 Co.sub.10 Fe.sub.10 
1010 
295 fcc-Al + Amo 
Duc Example 
69 Al.sub.45 Co.sub.35 Fe.sub.20 
-- 510 Com Bri Comparative example 
70 Al.sub.89 Co.sub.10.7 Fe.sub.0.3 
390 
105 fcc-Al + Amo 
Duc Comparative example 
71 Al.sub.98 Co.sub.1 Cu.sub.1 
320 
80 fcc-Al Duc Comparative example 
72 Al.sub.70 Co.sub.25 Cu.sub.5 
1005 
325 fcc-Al + Amo 
Duc Example 
73 Al.sub.45 Co.sub.35 Cu.sub.20 
-- 505 Com Bri Comparative example 
74 Al.sub..sub.89.7 Co.sub.10 Cu.sub.0.3 
485 
112 fcc-Al + Amo 
Duc Comparative example 
75 Al.sub.90 Co.sub.9 Mn.sub.0.5 Fe.sub.0.5 
996 
305 fcc-Al + Amo 
Duc Example 
76 Al.sub.89 Co.sub.8 Mn.sub.2 Cu.sub.1 
1210 
340 fcc-Al + Amo 
Duc Example 
77 Al.sub.90 Co.sub.7 Fe.sub.1 Cu.sub.1 
1005 
298 fcc-Al + Amo 
Duc Example 
78 Al.sub.90 Co.sub.7 Mn.sub.1 Fe.sub.1 Cu.sub.1 
1230 
310 fcc-Al + Amo 
Duc Example 
__________________________________________________________________________ 
TABLE 7 
__________________________________________________________________________ 
Alloy 
composition 
(Subscript numerals 
Sample 
represent atomic 
.sigma.f 
Hv Structural 
Bending 
No. percentage) 
(MPa) 
(DPN) 
state test 
__________________________________________________________________________ 
79 Al.sub.50 Fe.sub.40 Mn.sub.10 
-- 560 Com Bri Comparative example 
80 Al.sub.60 Fe.sub.35 Mn.sub.5 
845 
363 fcc-Al + Amo 
Duc Example 
81 Al.sub.65 Fe.sub.30 Mn.sub.5 
960 
375 fcc-Al + Amo 
Duc Example 
82 Al.sub.70 Fe.sub.20 Mn.sub.10 
875 
340 fcc-Al + Amo 
Duc Example 
83 Al.sub.85 Fe.sub.10 Mn.sub.5 
1070 
360 fcc-Al + Amo 
Duc Example 
84 Al.sub.95 Fe.sub.0.5 Mn.sub.4.5 
910 
260 fcc-Al + Amo 
Duc Example 
85 Al.sub.94 Fe.sub.0.3 Mn.sub.5.7 
480 
113 fcc-Al Duc Comparative example 
86 Al.sub.92 Fe.sub.6 Cu.sub.2 
1005 
276 fcc-Al + Amo 
Duc Example 
87 Al.sub.88 Fe.sub.8 Cu.sub.4 
1210 
302 fcc-Al + Amo 
Duc Example 
88 Al.sub.45 Fe.sub.35 Cu.sub.20 
-- 560 Com Bri Comparative example 
89 Al.sub.90 Fe.sub.6 Mn.sub.2 Cu.sub.2 
1112 
293 fcc-Al + Amo 
Duc Example 
90 Al.sub.75 Co.sub.8 Mn.sub.5 Ti.sub.12 
-- 511 fcc-Al + Com 
Bri Comparative example 
91 Al.sub.76 Fe.sub.4 Mn.sub.10 Ti.sub.10 
1210 
370 fcc-Al + Amo 
Duc Example 
92 Al.sub.78 Co.sub.4 Fe.sub.10 Zr.sub.8 
1100 
359 Amo Duc Example 
93 Al.sub.78 Fe.sub.8 Cu.sub.8 Ti.sub.6 
1060 
360 fcc-Al + Amo 
Duc Example 
94 Al.sub..sub.82 Co.sub.8 Mn.sub.3 Fe.sub.3 Zr.sub.4 
1090 
305 Amo Duc Example 
95 Al.sub..sub.83 Fe.sub.6 Mn.sub.3 Cu.sub.6 Ti.sub.2 
1206 
328 fcc-Al + Amo 
Duc Example 
96 Al.sub..sub.83 Co.sub.8 Mn.sub.4 Fe.sub.4 Zr.sub.1 
1230 
345 fcc-Al + Amo 
Duc Example 
97 Al.sub.98 Fe.sub.7 Cu.sub.4.5 Ti.sub.0.5 
1175 
339 fcc-Al + Amo 
Duc Example 
98 Al.sub.85 Fe.sub.10 Mn.sub.4.7 Zr.sub.0.3 
1049 
362 fcc-Al + Amo 
Duc Comparative example 
__________________________________________________________________________