U.S. patent application number 10/898240 was filed with the patent office on 2005-02-10 for annealing-induced extensive solid-state amorphization in metallic films.
This patent application is currently assigned to National Taiwan Ocean University. Invention is credited to Chu, Jinn P., Kuo, Chun-Hsin, Lo, Chang-Ting.
Application Number | 20050028900 10/898240 |
Document ID | / |
Family ID | 34076466 |
Filed Date | 2005-02-10 |
United States Patent
Application |
20050028900 |
Kind Code |
A1 |
Chu, Jinn P. ; et
al. |
February 10, 2005 |
Annealing-induced extensive solid-state amorphization in metallic
films
Abstract
An thin film alloy based on chemical elements with high glass
forming ability is disclosed. The alloy is deposited as a thin film
from a source of substantially the same chemical composition.
Within the deposited thin film, amorphization is induced
extensively up to decades of micrometers in size during controlled
annealing. Such controllable extensive amorphization throughout the
thin film is useful to regulate the proportion of amorphous phase
to crystalline phase, establish the structure/property
relationships and thus tailor specific properties.
Inventors: |
Chu, Jinn P.; (Keelung City,
TW) ; Kuo, Chun-Hsin; (Hsinchu City, TW) ; Lo,
Chang-Ting; (Taipei County, TW) |
Correspondence
Address: |
BACON & THOMAS, PLLC
625 SLATERS LANE
FOURTH FLOOR
ALEXANDRIA
VA
22314
|
Assignee: |
National Taiwan Ocean
University
Keelung
TW
|
Family ID: |
34076466 |
Appl. No.: |
10/898240 |
Filed: |
July 26, 2004 |
Current U.S.
Class: |
148/561 ;
148/403 |
Current CPC
Class: |
C22C 45/00 20130101 |
Class at
Publication: |
148/561 ;
148/403 |
International
Class: |
C22C 045/00 |
Foreign Application Data
Date |
Code |
Application Number |
Aug 4, 2003 |
TW |
92121265 |
Claims
What is claimed is:
1. An alloy film which is deposited as a film from a source
composed of desired chemical elements and annealed in a
controllable annealing process to form partly or fully amorphous
structures in the film, comprising a principal element with high
glass-forming ability and at least two secondary elements different
from said principal element; said principal elements with high
glass-forming ability are selected from a group consisting of iron,
cobalt, nickel, palladium, zirconium, titanium, magnesium, and
lanthanide series; and said secondary elements are selected from a
group consisting of aluminum, zirconium, copper, tin, zinc,
palladium, titanium, iron, cobalt, nickel, niobium, beryllium,
gallium, germanium, chromium, molybdenum, hafnium, lanthanide
series, VI.about.VIII group transition elements, phosphorus, boron,
carbon and silicon.
2. The alloy film of claim 1, wherein said partly amorphous
structure in the film refers to that there are
nanocrystallite/amorphous nanophase composite structures formed
extensively in the film.
3. The alloy film of claim 1, wherein said principal element with
high glass-forming ability is zirconium.
4. The alloy film of claim 3, wherein said secondary elements
comprise copper, aluminum and nickel.
5. The alloy film of claim 4, wherein the atomic percentage of said
alloy film is zirconium 40.about.60%, copper 15.about.35%, aluminum
5.about.20% and nickel 0.about.15%.
6. The alloy film of claim 5, wherein the atomic percentage of said
alloy film is zirconium 47%, copper 31%, aluminum 13%, and nickel
9%.
7. The alloy film of claim 1, wherein said principal element with
high glass-forming ability is iron.
8. The alloy film of claim 7, wherein said secondary elements
comprise cobalt, nickel, titanium, niobium and boron.
9. The alloy film of claim 8, wherein the atomic percentage of said
alloy film is iron 50.about.70%, cobalt 5.about.15%, nickel
5.about.15%, titanium 5.about.15%, niobium 0.about.10% and boron
0.about.20%.
10. The alloy film of claim 9, wherein the atomic percentage of
said alloy film is iron 65%, cobalt 8%, nickel 7%, titanium 13%,
niobium 1%, and boron 6%.
11. The alloy film of claim 1, wherein the thickness of said alloy
film is 0.2 .mu.m.about.50 .mu.m.
12. A process for manufacturing the alloy film of claim 1,
comprising the following steps: (a) using a alloy composed of
desired chemical elements as a film source; (b) depositing the
alloy onto a substrate to form a film; and (c) annealing the film
to induce partial or full amorphization in the film.
13. The process of claim 12, wherein the step of (b) is depositing
the alloy onto a substrate to form a film by physical vapor
deposition.
14. The process of claim 13, wherein said physical vapor deposition
is DC or RF magnetron sputtering.
15. The process of claim 12, wherein the step of (c) is annealing
the film by rapid thermal annealing to induce partial of full
amorphization in the film.
16. The process of claim 12, wherein the step of (c) is annealing
the film in an argon atmosphere.
17. The process of claim 12, wherein the heating rate of annealing
in step (c) is 5 K/min .about.200 K/min.
18. The process of claim 12, wherein the temperature of annealing
in step (c) is in the supercooled liquid region of the alloy
film.
19. The process of claim 12, wherein the temperature of annealing
in step (c) is in the range of 400 K.about.1200 K.
20. The process of claim 12, wherein the holding time of annealing
in step (c) is 10 seconds.about.3600 seconds.
Description
BACKGROUND OF THE INVENTION
[0001] 1. Field of the Invention
[0002] The present invention relates to an alloy film and a process
for manufacturing the same, especially to an annealing-induced
alloy film and a process for manufacturing the same.
[0003] 2. Description of the Related Art
[0004] Bulk metallic glass (BMG), also called bulk amorphous metal
(BAM) or bulk amorphous alloy (BAA), is a new material with
distinctive properties for special use. As compared with the
traditional amorphous metals formed from melting state rapidly,
this material has higher glass-forming ability (GFA). It can be
made into amorphous bulk form at extremely low cooling rates and
its thermal stability is better than that of the crystalline
metals. For example, A. L. Greer ("Metallic Glasses", Science, 267,
1947(1995)), A. Inoue ("Stabilization of Metallic Supercooled
Liquid and Bulk Amorphous Alloys", Acta Mater., 48, 279(2000)) and
Y Zhang, D. Q. Ahao, M. X. Pan and W. H. Wang ("Glass Forming
Properties of Zr-based Bulk Amorphous Alloys", J. Non-crystal.
Solids, 315, 206 (2003)) et al. summarized various GFA and BMG
systems, such as the ternary or quaternary alloys based on Fe, Co,
Ni, Pd, Zr, Ti, Mg, or La etc., or alloys with more components.
[0005] In addition to controlling the cooling rates, the
supercooled liquid region (.DELTA.T), defined by the difference
between the glass transition temperature (T.sub.g) and the
crystallization temperature (T.sub.x), can be broadened by varying
alloy components so that the superplasticity in this temperature
range can be utilized. Most BMG have good properties such as
mechanical properties, process ability of .DELTA.T, anti-corrosion,
hydrogen-stored ability, soft magnetism and other specific optical,
electrical, or chemical properties. The examples of actual
commercial applications of the material include exercise apparatus
and electrodes.
[0006] Crystallization generally tends to occur during thermal
annealing of amorphous alloy, which is characterized by the
distribution of various nano crystallites in the amorphous base and
found to strengthen the alloy structures. Researches on this
subject have been published by A. Inoue, C. Fan and a. Takeuchi
("High-strength Bulk Nanocrystalline alloys in a Zr-based System
containing Compound and Glassy Phases", J. Non-crystal. Solids,
250-252, 724(1999)) and A. L. Greer et al. ("Partially or Fully
Devitrified Alloys for Mechanical Properties", Mat. Sci. Eng. A,
A304-306, 68(2001)). Different synthesis processes would influence
the distribution, proportion, and routes of crystallization of the
nanocrystalites. In other words, if the relationships between
components, structures and properties of alloys were found, one
could synthesize the materials with desired properties by proper
synthesis process for specific applications.
[0007] In the solid state, amorphization could be achieved mainly
by mechanical alloying, solid-state amorphization (SSA), high
pressure and shock loading techniques. In all these solid-state
techniques except for the SSA, considerable energy is generally
required for ultimate amorphization. Metals can be also induced by
hydrogen to form amorphous hydride and proceed to powder metallurgy
for the solidification of BMG.
[0008] We hope that the characteristics of bulk amorphous metals
can be presented in the thin film, i.e. forming a metallic glass
thin film (MGTF) to extent the applications of those materials.
Similar to bulk alloys, the properties of MGTF can be modified by
controlling element components and nanocrystallites. The MGTF
presents good properties on mechanical isotropy, structure unity,
and have less crystal defects that the size effects should not be
existent, and its superplasticity as bulk alloys in the supercooled
liquid range is assistant to form three-dimension structures.
Therefore, MGTF can be applied in various fields, especially in the
fields of MEMS and optical record media.
[0009] Though the above-mentioned solid-state technique can
successfully synthesize bulk amorphous alloys, it is apparent that
the machining-involved technique is not suitable for films
formation and the use of SSA in film formations is limited. For
example, R. B. Schwarz and w. L. Johnson first proposed that the
solid-state amorphization in specific multilayer structures of
crystalline metals can be induced due to the annealing-induced
diffusion reactions ("formation of an amorphous alloy by
solid-state reaction of Pure Crystalline Metals", Phys. Rev. Lett.,
51, 415 (1983)). However, the extent of amorphization in this case
is trivial and confined to the reacted interface with the thickness
of few nanometers, as shown by B. X. Liu, W. S. Lai and z. J. Zhang
("Solid-State Crystal-to-Amorphous Transition in Metal-Metal
Multilayers and Its Thermodynamic and atomistic Modeling" Adv. In
Physics, 50, 367 (2001)).
[0010] Even various amorphous films can be formed through
traditional evaporation or sputtering process, the
annealing-induced crystallization reactions in both are quit
different and the properties of both are not certainly identical.
It is due to that the elemental components of traditional
deposition and sputter systems are not necessarily the same with
BMG and there are substantial differences between the film and bulk
manufacturing process. For example, Y Liu, S. Hata, K. Wada and A.
Shimokohbe et al. successfully sputtered the Pd-based ternary alloy
film on the aluminum layer or silicon wafer. They found that the
mechanical and thermal properties of obtained MGTF are similar to
those of BMG having the same components and those properties of the
MGTF are influenced by the sputtering conditions. However, they
didn't disclose the variance of crystallization, distribution or
properties caused by heating at different temperatures. They only
showed that, by way of the time-temperature-transition diagram, the
MGTF have good thermal stability while compared to the
crystallized, and observed that the resistance of the MGTF
decreased apparently when the annealing temperature increased to
T.sub.x.
[0011] S. Hata, Y Liu, T. Kato and A. Shimokohbe
("Three-dimensional Micro-Forming Process of Thin Film Metallic
Glass in the Supercooled Liquid Region", Proceeding of 10.sup.th
International Conference on Precision engineering (ICPE),
3741(2001)) sputtered the Zr-based ternary alloy films and
manufactured a three-dimensional cone spring by utilizing the
superplasticity of such films in the supercooled liquid temperature
region. According to the time-temperature-transition diagram, they
simply confirmed that heating the metals with a long time in the
temperature range would not result in crystallization, but they
didn't disclose the variance of crystallization, distribution or
properties caused by heating with different temperatures.
[0012] According to the above-mentioned prior art, a skilled person
in this field could not effectively control and anticipate the
structures and properties of BMG in the form of films. For example,
the distribution of crystalline phase/amorphous phase could not be
controlled extensively and uniformly. Thus, the requirement for
thin-film alloy utilization, not bulk alloys that have thicker
volumes, could not be satisfied yet.
SUMMARY OF THE INVENTION
[0013] Accordingly, the purpose of this invention is to produce an
alloy film based on elements with high glass-forming ability with
extensive amorphous structures and, in the meanwhile, effectively
control and anticipate its structures and properties so as to
satisfy the requirements mentioned above.
[0014] In one aspect, the present invention provides an alloy film
based on elements with high glass-forming ability, which is
deposited as a film from a source composed of desired chemical
elements, and annealed in a controllable annealing process to form
partly or fully amorphous structures in the film. The alloy film of
the invention includes a principal element with high glass-forming
ability and at least two secondary elements different from the
principal element, wherein the principal elements with high
glass-forming ability are selected from a group consisting of iron,
cobalt, nickel, palladium, zirconium, titanium, magnesium, and
lanthanide series; the secondary elements are selected from
aluminum, zirconium, copper, tin, zinc, palladium, titanium, iron,
cobalt, nickel, niobium, beryllium, gallium, germanium, chromium,
molybdenum, hafnium, lanthanide series, VI.about.VIII group
transition elements, phosphorus, boron, carbon, silicon or other
metal or nonmetal elements.
[0015] In a further aspect, the present invention provides a
process for manufacturing the above-mentioned alloy film,
comprising the following steps: using a alloy composed of desired
chemical elements as a film source, depositing the alloy onto a
substrate to form a film, and annealing the film to induce partial
or full amorphization in the film.
[0016] The present invention successfully induced partial or full
amorphization in the films, wherein the amorphous structures
distributed up to decades of micrometers and over the film
extensively. Such controllable extensive amorphization in the thin
film is useful to regulate the proportion of amorphous structure to
crystalline structure, establish the relationships between the
structure and property of the film and thus manufacture a film with
specific mechanical, electrical or optical properties.
BRIEF DESCRIPTION OF DRAWINGS
[0017] FIG. 1 shows (a) a differential scanning calorimetry (DSC)
thermogram of an as-deposited Zr-based film; variations of (b)
Knoop ultra-microhardness and (c) electrical resistivity with the
annealing temperature. Approximate location of supercooled liquid
region is marked by dash lines to facilitate visual comparison.
[0018] FIG. 2 shows the refractive index of the Zr-based films
annealed at different temperatures with different light
wavelengths. (a) refractive index v.s. wavelength; (b) refractive
index v.s. annealing temperature at 450 nm; (c) refractive index
v.s. annealing temperature at 500 nm; and (d) refractive index v.s.
annealing temperature at 550 nm.
[0019] FIG. 3 shows the refractive index of the Zr-based films
annealed at different temperatures with different light
wavelengths. (a) refractive index v.s. annealing temperature at 600
nm; (b) refractive index v.s. annealing temperature at 650 nm; (c)
refractive index v.s. annealing temperature at 700 nm; (d)
refractive index v.s. annealing temperature at 750 nm.
[0020] FIG. 4 shows the refractive index of the Zr-based films
annealed at different temperatures with different light
wavelengths. (a) refractive index v.s. annealing temperature at 800
nm; (b) refractive index v.s. annealing temperature at 850 nm; (c)
refractive index v.s. annealing temperature at 900 nm.
[0021] FIG. 5 shows plane-view transmission electron microscopy
(TEM) micrographs and diffraction patterns of the Zr-based films in
(a) as-deposited and annealed conditions at (b) 650 K, (c) 750 K,
(d) 800 K and (e) 850 K. The circled regions in (e) indicate the
locations for obtaining the diffraction patterns in (f).
[0022] FIG. 6 shows depth profiles of oxygen concentration in
uncapped Zr-based films in as-deposited and 800 K-annealed
conditions obtained by point analyses.
[0023] FIG. 7 shows the secondary ion mass spectrometry (SIMS)
depth profile of oxygen in (a) as-deposited and (b) 800 K-annealed
Zr-based films.
[0024] FIG. 8 shows a schematic illustration illustrating the free
energy relationship between the metastable sputtered phase (S),
amorphous phase (A), and stable crystalline phases (B and C) in the
Zr-based films annealed at low temperatures and 800 K. The arrows
indicate the film composition.
[0025] FIG. 9 shows electrical resistivity results of Fe-based
films in as-deposited and annealed conditions. DSC result is
included for comparison. Approximate location of supercooled liquid
region is marked by dash lines.
[0026] FIG. 10 shows TEM micrographs of the Fe-based films in (a)
as-deposited and annealed conditions at (b) 673 K, (c) 873 K and
(d) 923 K.
[0027] FIG. 11 shows TEM diffraction patterns of the Fe-based films
in (a) as-deposited and annealed conditions at (b) 673 K, (c) 873 K
and (d) 923 K.
[0028] FIG. 12 show the variation of in-plane (a) coercivity (Hc)
and (b) saturation magnetization with annealing temperatures of
Fe-based films were carried out by VSM. Approximate location of
supercooled liquid region is marked by dash lines.
[0029] FIG. 13 shows oxygen diffusion depth v.s. annealing
temperature plot for Fe-based films as measured by SIMS.
[0030] FIG. 14 shows a schematic illustration illustrating the free
energy relationship between the metastable phase (S), amorphous
phase (A), and stable crystalline phases (B and C) in the Fe-based
films annealed at low temperature and 923 K. The arrows indicate
the film composition.
[0031] FIG. 15 shows MFM images of Fe-based films at various
annealing temperatures. The black dots represent the attractive
force and the white dots represent the repulsive force between the
tip and sample
DETAILED DESCRIPTION OF THE INVENTION
[0032] The details of one or more embodiments of the invention are
set forth in the accompanying description below. Other features,
objects, and advantages of the invention will be apparent from the
detailed description, and the claims.
[0033] In one aspect, the present invention provides an alloy film
based on elements with high glass-forming ability, which is
deposited as a film from a source composed of desired chemical
elements, and annealed in a controllable annealing process to form
partly or fully amorphous structures in the film.
[0034] Generally, the usual BMG element systems are all suit for
the present invention. As known by those skilled in the art, the
principal element with high glass-forming ability includes iron,
cobalt, nickel, palladium, zirconium, titanium, magnesium, and
lanthanide series etc. Besides those principal elements, the
components of the alloy film include at least two secondary
elements that are different with the principal elements for
assisting in amorphization and forming ternary or quaternary alloys
or alloys with more components. The secondary elements include
aluminum, zirconium, copper, tin, zinc, palladium, titanium, iron,
cobalt, nickel, niobium, beryllium, gallium, germanium, chromium,
molybdenum, hafnium, lanthanide series, VI.about.VIII group
transition elements, phosphorus, boron, carbon, silicon or other
metal or nonmetal elements. It is apparent for one skilled person
to select and determine the proportion of the above-mentioned
elements. Thus the elements used in the invention are not limited
to the above-mentioned.
[0035] In one preferred embodiment of the invention, the principal
element with high glass-forming ability is zirconium. In this
embodiment, the secondary elements with high glass-forming ability
preferably include copper, aluminum and nickel. Take the
Zr--Cu--Al--Ni alloy for example, the atom percentage of the alloy
is zirconium 40.about.60%, copper 15.about.35%, aluminum
5.about.20% and nickel 0.about.15%; preferably, zirconium 47%,
copper 31%, aluminum 13% and nickel 9%, hereafter designated as
Zr.sub.47Cu.sub.31Al.sub.13Ni.sub.9
[0036] In the other preferred embodiment of the invention, the
principal element with high glass-forming ability is iron. In this
embodiment, the secondary elements with high glass-forming ability
preferably include cobalt, nickel, titanium, niobium and boron.
Take the Fe--Co--Ni--Ti--Nb--B alloy for example, the atom
percentage of the alloy is iron 50.about.70%, cobalt 5.about.15%,
nickel 5.about.15%, titanium 5.about.15%, niobium 0.about.10% and
boron 0.about.20%; preferably, iron 65%, cobalt 8%, nickel 7%,
titanium 13%, niobium 1% and boron 6%, hereafter designated as
Fe.sub.65CO.sub.8Ni.sub.7T.sub.13Nb.sub.1B.sub.6.
[0037] According to the invention, partly amorphous structures are
extensively formed in the film, i.e. there are extensive
nanocrystalite/amorphous nanophase composite structures formed in
the film. According to a preferred embodiment of the invention, for
example, the Zr.sub.47Cu.sub.31Al.sub.13Ni.sub.9 alloy film
annealed at 800 K, the amorphous structures are formed fully and
extensively in the film.
[0038] In one embodiment of the invention, the thickness of the
alloy film is about 0.2 .mu.m.about.50 .mu.m. In one preferred
embodiment of Zr-based alloy film, the thickness of the film is
about 5 .mu.m.about.110 .mu.m. In the other preferred embodiment of
Fe-based alloy film, the thickness of the film is about 0.5
.mu.m.about.10 .mu.m
[0039] In a further aspect, the present invention provides a
process for manufacturing the above-mentioned alloy film,
comprising the following steps: using a alloy composed of desired
chemical elements as a film source, depositing the alloy onto a
substrate to form a film, and annealing the film to extensively
induce partial or full amorphization in the film.
[0040] For depositing the thin film alloy, a proper film source (or
target) is prepared at first. For example, an alloy ingot with
desired chemical compositions is found by the vacuum arc remelting
process, and then proceed and heated into the final forms. A proper
substrate for deposition is selected according to the film types
and its application. The examples of substrates include
well-cleaned glass, silicon or other substrates made of different
materials.
[0041] The deposition process includes, for example, evaporation,
sputtering, or cathodic arc, preferably the physical vapor
deposition (PVD). The sputtering system used usually includes DC or
RF magnetron sputtering system. In one embodiment of the invention,
the working pressure and sputtering power during sputtering was
maintained at 3.times.10.sup.-3 torr and 100 W respectively, and
the sputtering system was in an argon atmosphere for sputtering.
The deposition conditions are apparent for one skilled person in
the art, thus the deposition conditions are determined in practice
without limited to the above-mentioned.
[0042] The as-deposited alloy film was then annealed with suitable
conditions to induce partial or extensive amorphization in the
film. Due to that the behavior of crystallization in the film is
influenced by the annealing conditions, the thermal behaviors of
those films composed of different elements during the heating
process is recorded by, for example, differential scanning
calorimeter (DSC) before annealing, establishing the temperature
range of supercooled liquid region and the corresponding variation
of film structures, whereby determining a preferred annealing
condition. The utilization of DSC results to determine annealing
conditions is also known by the skilled person.
[0043] Annealing can be conducted in any anneal furnace with
functions of adjusting parameters such as heating rates,
temperatures, etc. According to one embodiment of the invention,
the film was annealed in a rapid thermal annealing (RTA) system,
such as MILA-3000 RTA system. To prevent contamination, before
annealing, the annealing system was pumped down to 10.sup.-3 torr
range followed by purging with pure argon, such as 99.9995% Ar, for
several times to minimize the residual reactive gases, such as
oxygen.
[0044] The heating rate of annealing is about 5 K/min.about.200
K/min. In a preferred embodiment of Zr-based alloy film, the
heating rate of annealing is 40 K/min. In the other preferred
embodiment of Fe-based alloy film, the heating rate of annealing is
100 K/min. The temperature of annealing is preferably on the range
of supercooled liquid region in order to control the reaction rate
of amorphization. The supercooled liquid region varies with the
components of alloys. For example, the temperature range of
annealing is about 400 K.about.1200 K. In a preferred embodiment of
Zr-based alloy film, the temperature range of annealing is about
550 K .about.950 K. In the other preferred embodiment of Fe-based
alloy film, the temperature range of annealing is about 673
K.about.1073 K. The holding time of annealing is about 10
s.about.3600 s, preferably 60 s. The cooling rate of annealing is
about 5 K/min.about.200 K/min, preferably 20 K/min.about.40
K/min.
EXAMPLE 1
Zr-based thin film
[0045] Experimental Procedure
[0046] Zr-based quaternary alloy thin films of thickness 5-10 .mu.m
with a nominal composition of Zr.sub.47Cu.sub.31Al.sub.13Ni.sub.9
were deposited onto the well-cleaned glass substrate using an RF
magnetron sputtering system in an argon atmosphere. The working
pressure and RF power during sputtering were maintained at
3.times.10.sup.-3 torr and 100 W, respectively. The compositions of
the films were measured using an electron probe for microanalysis
(EPMA). The compositional fluctuation at various points on the film
surface was also determined and found to be very small, (around 1%)
which reveals the uniformity of the deposited films. The films were
then annealed in a rapid thermal annealing (RTA) system in Ar at
temperatures ranging from 550 to 950 K. To avoid contamination, the
RTA system was pumped down to 10.sup.-3 torr range followed by
purging with pure Ar for several times. For RTA, the samples were
kept 60 seconds in holding time with the heating rate of 40 K/min.
The crystallization of the film was studied using a differential
scanning calorimeter each film was delaminated entirely from the
glass substrate for the DSC analysis without the aid of any
chemical solutions. The crystal structures of films were examined
by a transmission electron microscopy (TEM). Broad-face TEM sample
discs were thinned from the substrate side by a dimpler, followed
by an ion miller for the final perforation. TEM examinations were
performed at 300 keV. Sheet resistance measurements of films were
done at room temperature by a four-point probe method. For the
mechanical property evaluation, the Knoop ultramicrohardness of
film was measured. To eliminate any error due to the substrate
effect, the indentation was applied with 25 g loading, 15 second
holding time and loading rate of 40 .mu.m/sec.
[0047] Results and Discussion
[0048] The DSC study in FIG. 1(a) reveals the as-deposited film
undergoes a glass transition (Tg) and crystallization (Tx) at 758.3
K and 797.2 K, respectively. The supercooled liquid region
(.DELTA.T), conventionally defined by the difference between Tg and
Tx, is thus measured to be .about.39 K. This value is comparable to
32 K reported for sputtered Pd.sub.76 Si.sub.17Cu.sub.7 film, but
smaller than 70 K for Zr.sub.75Cu.sub.19A.sub.l6 film. However, the
.DELTA.T values in the thin film form are generally lower than in
bulk form. .DELTA.T values are 98 K and 81 K for
Zr.sub.55Cu.sub.30Al.sub.10Ni.sub.5 and Pd.sub.40Ni.sub.40P.sub.20
bulk metallic glasses, respectively. The differences in chemical
composition and preparation methods may be attributed to the
difference in .DELTA.T. Annealing of the film at low temperatures
results in the release of residual stress present in the
as-deposited condition and thus causes the decrease in film
hardness, as seen in FIG. 1(b). When the film is annealed below the
glass transition temperature, the film hardness gradually increases
with the annealing temperature and reaches a hardness value of
.about.HK670 at 750 K. Such a beneficial strengthening effect is
presumably attributed to the nanocrystallite/amorphous composite
structure (or called nanophase composite). The film hardness then
decreases slightly to .about.HK630 at 800 K for annealing
temperatures within the supercooled liquid region, indicating a
possible structure change in this temperature range. The hardness
increases noticeably for the films annealed at temperatures above
the crystallization transition, followed by a decrease in hardness
due to an extensive crystallization and grain growth. With the
exception of the anomalous drop in the .DELTA.T region, the film
hardness in general follows an increasing trend with annealing
temperature, reaching a maximum at 900 K with a .about.60% increase
from the as-deposited value. While most of other sputtered films
showing a negative temperature dependence of film hardness, the
films prepared in the present example revealed an increasing
hardness property with annealing temperature. This distinctive
behavior may be a characteristic of annealed metallic glass thin
films with the nanophase composite structure, which could lead to a
new class of high performance nano materials. For the electrical
resistivity result in FIG. 1(c), the annealing in the AT region
exhibits a different behavior. An abrupt increase (.about.40%) in
film resistivity to .about.65 .mu..OMEGA.-cm at 800 K is noted for
the film annealed in the .DELTA.T region. This again suggests there
might be a distinct structure in the supercooled liquid region.
Yet, annealing at temperatures below the glass transition results
in a decrease in the film resistivity, to .about.46 .mu..OMEGA.-cm
at 750 K. Such a decrease is originated from the combined effects
of stress release, nanocrystallite growth, and reshuffling of
sputtered structure, analogous to those of other sputtered metallic
films. In general, a decreasing trend of the film resistivity with
the annealing temperature is observed except for a dramatic
increase in the supercooled liquid region. When compared with that
of the sputtered Pd-based glass-forming alloy film, a similar
increase in the overall decreasing trend of film resistivity near
the glass transition has been reported by Liu et al. However, due
to the lack of crystal structure/microstructure details and no
resistivity values reported in the supercooled liquid region, it is
inconclusive to determine whether full amorphization in their films
is obtained in the supercooled liquid region.
[0049] Except varying with temperature and light wavelengths, the
refractive index varies with elemental components of the material.
The refractive indexes of the films annealed at different
temperatures v.s. light wavelengths are show in FIGS. 2, 3 and 4.
When the film is annealed at temperatures below the glass
transition temperature, such as a temperature below 750 K, the film
refractive index decreases for all given wavelengths. Such
phenomenon is presumably attributed to the nanophase composite. The
film refractive index then slightly increases about 0.2 with
annealing temperatures in the supercooled liquid region, such as
750.about.800 K, indicating that the film structure changed in this
temperature range. The film refractive index decreases again when
the annealing temperature increased over the crystallization
temperature, such as over than 800 K, due to the occurrence of
crystallization. For example, the minimum refractive index occurs
at 825 K. Then with temperatures increasing again, the crystal
growth is more remarkable and the refractive index increases, while
the increasing range and the given wavelength shows an inverse
relationship. For example, the film refractive index for short
wavelength light may increase back to that of as-deposited film,
but the increase of refractive index for long wavelength light is
unapparent. Briefly, except for the supercooled liquid region, the
film refractive index substantially decreases with increasing
temperatures, but it changes with light wavelength when annealed at
a temperature above the crystallization temperature.
[0050] FIG. 5 represents the TEM results of the films in the
as-deposited and annealed conditions. The TEM image in FIG. 5(a)
and a typical corresponding diffraction pattern indicate the
dominant structure in the as-deposited condition is the
nanocrystalline structure and the amorphous is found to be a minor
phase. The nanocrystallites have sizes in a range of 10-30 nm and
thus produce a spotty-like ring diffraction pattern as shown in
FIG. 5(a). Measurements performed on the diffraction pattern show
that the structure is either tetragonal Zr.sub.2Ni (JCPDS 38-1170)
or cubic Zr.sub.2Ni (JCPDS 41-0898). After annealing at 650 K, the
thermal-activated stress release and reshuffling of sputtered
structure promote better amorphization and refined crystallites,
developing a glassy (amorphous) matrix containing uniformly
dispersed nanocrystalline phases, as indicated in the TEM image in
FIG. 5(b). The refined crystallites reveal sizes in the range
between .about.10 and 20 nm. The well-defined, continuous ring
diffraction pattern in FIG. 5(b) suggests the film consists of the
rather fine nanocrystalline structure than that in the as-deposited
condition. When annealed at 750 K, about the onset of the glass
transition, the agglomeration and growth of nanocrystalline phases
in the form of crystalline networks embedded in the amorphous
matrix are seen in FIG. 5(c). More crystalline spots in the
diffraction pattern of FIG. 5(c) indicate these phases are better
crystallized. The crystal structure, however, could not be
unambiguously identified.
[0051] Annealing at 800 K, in the .DELTA.T region, produces a fully
amorphous structure without any observable crystalline phases, as
evidenced by the TEM image of FIG. 5(d) and by the typical broad
diffuse diffraction rings characteristic of a glass in FIG. 5(d).
The presence of this distinct amorphous structure thus explains
anomalous properties observed in the .DELTA.T region in FIGS. 1(b)
and (c). For the 850 K-annealed films, above the crystallization
transition, the crystalline phases increase in distribution and
grow in size compared to the sample annealed at 750 K (FIG. 5(e))
as a result of devitrification. The crystallinity of the glassy
matrix also improved considerably as evidenced by the spotted rings
in the diffraction pattern in FIG. 5(e). The crystalline phases
could not be definitely identified; but they are likely to be
Zr.sub.2Ni. After annealing at 950 K, well above the
crystallization transition, the film develops a homogenous
structure consisting of a uniformly dispersed large crystalline
phase in a nanocrystalline matrix (FIG. 5(e)). The structure of the
large crystalline phase is the same as that observed in the sample
annealed at 850 K. Annealing at temperatures well above the
crystallization transition results in an extensive growth of
crystallites, thus deteriorating the film hardness and resistivity
properties shown in FIGS. 1(b) and (c).
[0052] We now discuss the mechanism by which the partly amorphous
film dispersed with crystallites is transformed to the fully glassy
state during annealing. Since impurities such as hydrogen have been
reported to induce amorphization in many alloys under highly
pressurized (.about.5 MPa) and thermal environment, species in the
substrate and residual gases in the RTA system might have possibly
played roles to promote the amorphization in our case. However,
such possibilities are not likely because the short annealing
duration (60 seconds) and relatively low operational pressure are
unfavorable to allow the diffusion-driven reactions to occur.
Furthermore, oxygen impurity and contamination are known to
strongly affect the stability of the amorphous phase in BMG. As Zr
and its alloys are known to be susceptible to oxidation,
incorporation of oxygen into the film is often very difficult to
avoid even in the reduced and protective annealing environment.
However, the oxygen effect would be negligible for the annealing
temperatures .ltoreq.800 K on account of the following reasons.
Interaction of Zr with oxygen involves both oxygen dissolution and
formation of scale (mainly ZrO.sub.2). The scale formation was
limited in this study because our X-ray diffraction analysis
results (which are not shown here) revealed the presence of oxide
phases only at and above 850 K. The oxygen dissolution during
annealing is thus considered in this study. In the absence of an
exact diffusion coefficient, D, we take an Arrhenius expression,
D(in cm.sup.2/sec)=5.2exp [-212/(8.314*T)], where T is absolute
temperature. This expression has been commonly used for Zr and its
alloys to determine the depth of oxygen dissolved into the metal
through volume diffusion. Based on X=2{square root}{square root
over (Dt)}, where t is the annealing time and X the distance at
which the oxygen concentration falls to half the initial value of
maximum solubility (28.5%) at the metal oxide interface, and an
assumption that the time spent in heating up to 0.8 of annealing
temperature makes insignificant contributions to the total amount
of diffusion, depths of oxygen dissolution into our annealed films
are estimated to be 0.1, 0.3, 0.7, 1.4 .mu.m for 800, 850, 900 and
950 K, respectively. The estimate is consistent with the depth
profile of oxygen concentration shown in FIG. 6. This figure
indicates no oxygen contamination due to annealing at 800 K. The
average oxygen content in the as-deposited film is, within the
experimental error, nearly equal to that of the sample annealed at
800 K. The oxygen content of the Zr-based film is performed by
secondary ion mass spectrometry (SIMS), as shown in FIG. 7. The
depth profile of oxygen shows that the oxygen dissolution in the
as-deposited film is 20 nm, while in the 800 K-annealed film is 80
nm. 80 nm in depth is considered to be insignificant for the whole
film thickness of 10 .mu.m. Based on FIGS. 6 and 7, we conclude
that the formation of amorphous phase during the 800 K annealing is
not associated with oxygen contamination at all. Furthermore, no
apparent thickness dependence of amorphization for the 5-10 .mu.m
film thickness range examined suggests the negligible effect of
this diffusion-driven event.
[0053] To interpret the present results and explain the
thermodynamics that govern the structure evolution during
annealing, hypothetical free-energy diagrams are shown in FIG. 8.
The diagrams, though approximate, contain the essential features
required to interpret the present results. The relative position of
the free energy curves in the figure is rationalized based on
thermodynamic considerations. In this figure, of primary importance
is the large negative heat of mixing observed for the Zr-based
alloy system. Inoue has suggested that the large negative heat of
mixing enables multicomponent alloys to readily form a glassy phase
during cooling from the melt. The large negative heat of mixing and
hence the low free energy for the amorphous phase can be related to
the amorphization in our case. The low free energy serves as the
thermodynamic driving force for the amorphous phase to form as an
intermediate phase during annealing, which eventually becomes
thermodynamically stable by full crystallization at high
temperatures. Based on the schematic, it is anticipated that
similar amorphization behavior could take place in other sputtered
glass-forming films with large negative heat formation.
[0054] It has been proposed that sufficient thermal-driven
diffusions and high interfacial energies between two different
phases are prerequisite for the SSA to take place. Our result thus
demonstrates, as we qualitatively hypothesized, the thermal energy
supplied by the annealing and the interfacial energy arising from
the nanocrystallite/glass interfaces might have imparted
significant influences on the amorphization and crystallization of
film during annealing. Upon annealing at below the supercooled
liquid region, the interfacial energy appears to be predominant
because of the presence of nanocrystallites, and hence the growth
of nanocrystalline phases is favorable. In addition, the viscosity
of glass matrix presumably decreases with the annealing
temperature, analogous to that required for the superplasticity
behavior of metallic glasses. The glassy matrix with sufficiently
low viscosity thus allows the nanocrystalline phases to "move
around" in the matrix. Ultimately, nanocrystallites become
interconnected and form a network structure. Once this
interconnected crystalline network structure is formed,
facilitating the flow of electrical current, a decrease in
resistivity is thus expected (FIG. 1(c)). The thermal energy in
this case, however, is insufficient due to the low annealing
temperatures and subsequently the amorphization occurs with limited
extents, in spite of the relatively high interfacial energy. Yet,
the extent of amorphization increases resulting from sufficient
thermal-driven diffusions for the high annealing temperatures.
Eventually, for the full amorphization at 800 K, as-sputtered
nanocrystallites are thermally annihilated and "liquefied" into the
glassy matrix, on account of the combined effects of sufficient
thermal energy and excessive interfacial energies. Contrary to the
planar interfacial area in SSA multilayer films, our irregular and
plentiful nanocrystallite/glass interfaces that are present
throughout the film are considered to yield much higher interfacial
energies. This, in turn, leads to the possibility of the occurrence
of large-scale amorphization in our film. While the extensive
amorphization reported in this study is beneficial for ease of
characterizations and readiness for potential applications,
evaluations of additional properties (such as optical) and
searching for other possible glass-forming systems exhibiting this
phenomenon will be the subjects of further work. Moreover, the
controllable amorphous structure may serve as a precursor for a new
class of nano materials.
[0055] In summary, here we show direct experimental evidences that
annealing of Zr.sub.47Cu.sub.31Al.sub.13Ni.sub.9 film at a
temperature within the supercooled liquid region results in
extensive amorphization, presumably attributed to sufficient
thermal and interfacial energies between nanocrystallites and
glassy matrix. The formation of comprehensive amorphous structure
gives rise to notable alterations in the electrical and mechanical
properties of annealed film. Additional features of the present
work are that a prominent strengthening effect is observed due to
the improved amorphous matrix dispersed with nanocrystalline phases
upon annealing and that one can take this advantage to tailor the
film properties by modulating the amorphous content in the annealed
films. The controllable amorphization may also serve as a precursor
for exciting new nano materials.
[0056] The details of above-mentioned Zr-based alloy film have been
published on Physical Review B, Vol. 69, page 113410 in March of
2004.
EXAMPLE 2
Fe-based thin films
[0057] Experimental procedure
[0058] The Fe.sub.65CO.sub.8Ni.sub.7T.sub.13Nb.sub.1B.sub.6 (atomic
percent, at. %) thin films were prepared by anRF magnetron
sputtering method. The Fe-Co-Ni-Ti-Nb-B target was an as-cast
alloy. Thin films of thickness 0.5-10 .mu.m were deposited on glass
substrate. The deposition was carried out under the following
conditions. The base vacuum was 10.sup.-7 Torr, Ar gas flow rate
was 20 sccm, and the working pressure was 3.times.10.sup.-3 Torr.
The power of 100 W was applied during the deposition. The film was
annealed in Ar at a heating rate of 100 K/min and a holding time of
60 s at temperatures ranging from 673 to 1073 K. The annealing
system was pumped down to the 10.sup.-3 Torr range followed by
several purging with Ar. Compositions of thin films were determined
by Electron Probe Microanalyzer. The thermal behavior of the film
was determined using a differential scanning calorimeter (DSC) in
Ar at a scanning rate of 100 K/min. The DSC film sample was
delaminated from the glass without the aid of any chemical
solutions. Sheet resistance and hardness measurements were carried
out by four-point probe and Knoop ultramicrohardness methods,
respectively. The ultramicrohardness was measured with a 25 g
loading, a 15 s holding time, and a loading rate of 40 .mu.m/s. The
crystal structures of films were examined by a transmission
electron microscopy (TEM). TEM examinations were performed at 200
keV. The composition distribution and oxygen content of the films
were performed by secondary ion mass spectrometry (SIMS, Cameca
IMS6F), depth-profiling studies were carried out by Ar sputtering.
The coercive field (Hc) and the saturation magnetization were
measured at room temperature by a vibrating sample magnetometer
(VSM), using maximum field strength of 7500 Oe.
[0059] Results and Discussion
[0060] FIG. 9(a) shows the DSC scan of as-deposited thin film at
scan rate of 100 K/min. From FIG. 9(a), the Fe-based thin film
metallic glass exhibits a slope change due to glass transition at
about 816 K and exothermic reactions due to crystallization at
about 866 K and 950 K. Thus, its glass transition temperature
(T.sub.g), crystallization temperature (T.sub.x) are 816 K and 866
K, respectively. The supercooled liquid region (.DELTA.T), defined
as the difference between T.sub.g and T.sub.x, is thus measured to
be .about.50 K. When compared this value with the reported for
Fe--Co--Ni--Ti--Nb--B and Fe--Al--Ga--P--C--B alloys, our .DELTA.T
value was smaller than those studies, even though there were
differences in the apparatus and heating rate conditions between
this work and those studies. Based on the DSC results, the
annealing temperatures of the films were determined. The DSC
results also reveal the heat of exothermic reaction (.DELTA.H) of
thin film from the area under exothermic peaks. The two exothermic
peaks are considered be caused by the nucleation (the first peak)
and by the growth of nuclei (the second peak), as proposed in other
amorphous thin films. Table 1 lists summary results of DSC for
as-deposited film at scan rate of 100 K/min.
1 TABLE 1 Crystallization Temperature (K) Scan Rate Glass
Transition Onset Peak .DELTA.H (K/min) Temperature (K) (K) (K)
(J/g) .DELTA.T (K) Fe-based 100 816 866 913 765 50 thin film
[0061] Thermal annealing results in variations of
microstructure/crystal structure in the film and thus film
properties such as hardness and electrical resistivity have been
altered. Annealing of the Fe-based films at low temperatures yields
the release in the residual stress present in the as-deposited
condition and thus causes the decrease in film hardness, as seen in
FIG. 9(b). When the films are annealed at temperatures below the
glass transition, the film hardness gradually increases with the
annealing temperature, reaching .about.HK967 at 773 K. Such
beneficial strengthening effect is presumably attributed to the
nanocrystallite/amorphous composite structure (or called nanophase
composite). Similar hardness variation of film due to this
nanophase composite has been also reported in the previous example
of the Zr-based film during annealing. The film hardness decreases
slightly to .about.HK867 at 823 K, and .about.HK1048 at 973 K,
indicating a possible structure change in this temperature range.
The hardness increases for the films annealed at temperatures above
the crystallization transition, followed by a decrease in hardness,
as a result of extensive crystallization and grain growth. With
exception of the anomalous drop in the .DELTA.T region, the film
hardness in general follows an increasing trend with annealing
temperature, reaching a maximum at 1023 K with a .about.42%
increase from the as-deposited value. While most of other sputtered
films showing a negative temperature dependence of film hardness,
the film prepared in the present study has a hardness property
increased with annealing temperature. When compared with that of
the sputtered Zr-based glass-forming alloy film, the Fe-based film
in the present study reveals similar increasing trend of hardness
variation with annealing temperature except for the two decreases
at 823 K and 973 K. The distinctive behavior of positive
temperature dependence of hardness may be a characteristic of
annealed glass-forming thin films with the nanophase composite
structure, which could lead to a new class of high performance
materials.
[0062] FIG. 9(c) shows the electrical resistivity results of
Fe-based films in as-deposited and annealed conditions. From this
figure, the annealing in the .DELTA.T region causes different
behavior. It is noted an abrupt increase in film resistivity to
.about.255 .mu..OMEGA.-cm at 848 K, then decreased to .about.233
.mu..OMEGA.-cm at 873 K for the film annealed in the .DELTA.T
region. This again agrees with the distinct structure change in the
supercooled liquid region. Yet, annealing at temperatures below the
glass transition results in an abrupt decrease in the film
resistivity to .about.224 .mu..OMEGA.-cm at 773 K, then slightly
increased to .about.228 .mu..OMEGA.-cm at 798 K. Such a transition
is originated from the combined effects of stress release,
nanocrystallite growth, and reshuffling of sputtered structure,
analogous to those of other sputtered metallic films. In general, a
decreasing trend of our film resistivity with the annealing
temperature is observed except for a dramatic transition in the
supercooled liquid region.
[0063] FIGS. 10 and 11 show TEM bright-field images and
corresponding diffraction patterns, respectively, of Fe-based films
in as-deposited and annealed conditions. The as-deposited film
exhibits nanocrystalline peaks dispersed in an amorphous matrix.
The sizes of nanocrystalline phases range from 2 nm to 5 nm.
According to the d-spacing measured from the diffraction pattern
(FIG. 11 (b)), the nanocrystalline are likely to be cubic FeNi
(JCPD#18-0645). The annealing at 673 K for 60 s yields better
crystallinity and growth of nanocrystalline phases with amorphous
structure becoming indistinct. These are evidenced by twinned
grains in TEM images (FIG. 10(b)) and well defined spotty rings in
diffraction patterns (FIG. 11(b)). Based on FIG. 11(b), the
d-spacing of diffraction rings are 2.048 .ANG., 1.79 .ANG., 1.27
.ANG., 1.08 .ANG. and 0.81 .ANG., respectively. The sizes of the
nanocrystalline are in a range of 5 to 9 nm. The diffused spotty
rings in diffraction patterns (FIG. 11(c)), and poorly defined
crystalline phases in the amorphous matrix (FIG. 10(c)) at 873 K
indicate that the nanocrystalline phases appear to dissolve in the
amorphous matrix and lose their crystallinity. With 50 K increase
in annealing temperature to 923 K, the film reveals more amorphous
structure in the matrix with less nanocrystalline phase. The TEM
results shown here are somehow consistent with those reported
earlier in Zr-based glass-forming films.
[0064] FIG. 12(a) shows the variation of in-plane coercivity (Hc)
and saturation magnetization with various Fe-based films in
as-deposited and annealed conditions. As shown in FIG. 12(a), when
the films are annealed at temperatures below the glass transition,
the coercivity increases with the annealing temperature, reaching
97.5 Oe at 798 K. The film coercivity gradually decreased to 33 Oe
at 873 K for annealing temperatures within the supercooled liquid
region, and then the film coercivity increased to 114 Oe at 898 K.
The coercivity gradually decreased for the films annealed at
temperatures above the crystallization transition, followed by a
significant increase in coercivity, which corresponds to extensive
crystallization and grain growth. Thus, cyclic variations of
coercivity and saturated magnetization with annealing temperature
are observed. A detailed analysis is needed in order to understand
such cyclic behavior of magnetic properties. From FIG. 12(b), we
also find the cyclic behavior of saturated magnetization. The
saturated magnetization slightly decreased to 0.0112 emu/g at 673
K. When the annealed condition at 798 K, the saturated
magnetization increased to 0.0152 emu/g. The minimum saturated
magnetization was 0.0099 emu/g at 823 K, and the maximum saturated
magnetization was 0.0271 emu/g at 848 K, for annealing temperatures
within the supercooled liquid region. The saturated magnetization
gradually dropped to 0.013 emu/g at 948 K, and then up to 0.0192
emu/g at5 973 K for annealing temperature above crystallization
temperature. Based on FIG. 12(a) and (b), we found that the low
coercivity (.about.33 Oe) and maximum saturated magnetization
(.about.0.0271 emu/g) within supercooled liquid region. Further
studies on such cyclic behavior at different annealing temperatures
are needed in order to establish better understanding on the
magnetic of this film.
[0065] Since impurities such as hydrogen have been reported to
induce amorphization in many alloys as mentioned in the previous
example of Zr-based film, FIG. 13 is the oxygen diffusion depth of
the Fe-based film v.s. annealing temperature plot, as measured by
SIMS. According to this figure, the oxygen diffusion depth up to
130 nm is insignificant when compared with the whole film thickness
of 10 .mu.m. Thus, it is concluded that the oxygen impurity effect
is negligible.
[0066] In this example, qualitatively, the thermal energy supplied
by the annealing and the interfacial energy arising from the
nanocrystalline/matrix interfaces have significant influences on
the amorphization and crystallization of film during annealing.
Upon annealing, our films clearly show a structure development
sequence of metastable sputtered structure .fwdarw.metastable
nanocrystallite/amorpho- us nanophase composite .fwdarw.single
metastable amorphous phase .fwdarw.stable crystalline structure. To
interpret the present results and explain the thermodynamics that
govern the structure evolution during annealing, hypothetical free
energy diagrams are shown in FIG. 14. The diagrams, though
approximate, contain the essential features required to interpret
the present results. The relative position of the free energy
curves in the figure is rationalized based on thermodynamic
considerations. In this figure, of primary importance is the large
negative heat of mixing observed for the Fe-based alloy system. As
stated in the previous section, has suggested that the large
negative heat of mixing enables multicomponent alloys to readily
form a glassy phase during cooling from the melt. The large
negative heat of mixing and hence the low free energy of the
amorphous phase can be related to the amorphization in our case.
The low free energy provides the thermodynamic driving force for
the amorphous phase to form as an intermediate phase during
annealing, which eventually becomes thermodynamically stable by
full crystallization at high temperatures.
[0067] In addition to the thermodynamic factor, the thermal energy
at elevated temperatures and the interfacial energy arising from
the nanocrystallite/glass interfaces are kinetically favorable for
the amorphization, as proposed previously in SSA. At low
temperatures, the large interfacial energy drives coarsening of the
metastable crystalline phase through a process analogous to Ostwald
ripening. Further, the viscosity of the glass matrix presumably
decreases with temperature, analogous to that required for the
superplasticity behavior of metallic glasses. The amorphous matrix,
with a sufficiently low viscosity, allows the metastable
nanocrystalline phases to "move around" and reorient in the matrix.
An extensive amorphization occurs as the metastable
nanocrystallites are thermally annihilated and "liquefied" into the
amorphous matrix due to the combined effects of sufficient thermal
energy and excessive interfacial energy. In contrast to the planar
interfacial area in SSA multilayer films, the single layer films in
this study are considered to yield much higher interfacial energies
since nanocrystallite/amorphous matrix interfaces are present
through the film thickness. As a result, wide spread amorphization
could take place in the film. Above T.sub.x, the amorphous
structure is no longer thermodynamically favorable and
crystallization readily proceeds.
[0068] Conclusions
[0069] In this example, the Fe-based film shows the extensive
amorphization phenomenon, presumably attributed to sufficient
thermal and interfacial energies between nanocrystallites and
amorphous matrix. The extensive amorphization gives rise to
distinct variation in the electrical, mechanical and magnetic
properties of annealed Fe-based film within supercooled liquid
region. However, Fe-based film did not form fully amorphous
structure as Zr-based film did. This might be due to the fact that
Zr-based film has better glass-forming ability than Fe-based film.
Some important results are summarized as follows.
[0070] (1) The electrical resistivity result indicates that
Fe-based film has the high resistivity (.about.228 .mu..OMEGA.-cm)
within the supercooled liquid region.
[0071] (2) The Fe-based film has low ultra-microhardness
(.about.866.6 HK) with in the supercooled liquid region.
[0072] (3) TEM results show the Fe-based film has the amorphization
taking place at 923 K.
[0073] (4) The VSM result shows the Fe-based film has low
coercivity (.about.33 Oe) and maximum saturated magnetization
(.about.0.027 emu/g) within the supercooled liquid region. Cyclic
variation of coercivity and saturated magnetization with annealing
temperature are observed.
[0074] (5) The MFM result (FIG. 15) shows the domain structure of
Fe-based film in agreement with electrical resistivity and
TEMresults.
[0075] OTHER EMBODIMENTS
[0076] All of the features disclosed in this specification may be
combined in any combination. Each feature disclosed in this
specification may be replaced by an alternative feature serving the
same, equivalent, or similar purpose. Thus, unless expressly stated
otherwise, each feature disclosed is only an example of a generic
series of equivalent or similar features.
[0077] From the above description, one skilled in the art can
easily ascertain the essential characteristics of the present
invention, and without departing from the spirit and scope thereof,
can make various changes and modifications of the invention to
adapt it to various usages and conditions. Thus, other embodiments
are also within the scope of the following claims.
* * * * *