U.S. patent application number 12/933825 was filed with the patent office on 2011-01-27 for alloy and method for producing alloy.
This patent application is currently assigned to CANON KABUSHIKI KAISHA. Invention is credited to Toshiya Inoue, Jun Tamai.
Application Number | 20110017365 12/933825 |
Document ID | / |
Family ID | 40757019 |
Filed Date | 2011-01-27 |
United States Patent
Application |
20110017365 |
Kind Code |
A1 |
Tamai; Jun ; et al. |
January 27, 2011 |
ALLOY AND METHOD FOR PRODUCING ALLOY
Abstract
To provide an alloy which can suppress minute temporal
deformation of a Super Invar alloy as much as possible, and a
method for producing the alloy. The alloy of the present invention
includes iron, nickel, and cobalt, which are the basic components
of a Super Invar alloy, and is characterized in that an amount of a
fraction which has not carbidized in carbon contained in the alloy
is 0.010 wt % or less.
Inventors: |
Tamai; Jun; (Yokohama-shi,
JP) ; Inoue; Toshiya; (Utsunomiya-shi, JP) |
Correspondence
Address: |
FITZPATRICK CELLA HARPER & SCINTO
1290 Avenue of the Americas
NEW YORK
NY
10104-3800
US
|
Assignee: |
CANON KABUSHIKI KAISHA
Tokyo
JP
|
Family ID: |
40757019 |
Appl. No.: |
12/933825 |
Filed: |
April 23, 2009 |
PCT Filed: |
April 23, 2009 |
PCT NO: |
PCT/JP2009/058535 |
371 Date: |
September 21, 2010 |
Current U.S.
Class: |
148/547 ;
148/328 |
Current CPC
Class: |
C22C 38/08 20130101;
C21D 6/001 20130101; C21D 2211/004 20130101; C21D 6/02 20130101;
C21D 8/005 20130101 |
Class at
Publication: |
148/547 ;
148/328 |
International
Class: |
C21D 8/00 20060101
C21D008/00; C22C 38/20 20060101 C22C038/20; C21D 6/02 20060101
C21D006/02; C22C 38/08 20060101 C22C038/08 |
Foreign Application Data
Date |
Code |
Application Number |
Apr 28, 2008 |
JP |
2008-117352 |
Apr 13, 2009 |
JP |
2009-097226 |
Claims
1. An alloy comprising iron, nickel, and cobalt which are basic
components of a Super Invar alloy, characterized in that an amount
of a fraction which has not been carbidized in carbon contained in
the alloy is 0.010 wt % or less.
2. The alloy according to claim 1, characterized by further
comprising carbide forming elements, wherein carbides formed from
at least a part of the carbide forming elements and carbon are
dispersed in a base phase.
3. The alloy according to claim 2, characterized in that a total
carbon content in the alloy is 0.010 wt % or more.
4. The alloy according to claim 3, characterized in that the
carbide forming elements are contained in a higher amount than the
total carbon content in terms of atomic percent.
5. The alloy according to claim 2, characterized in that a content
of the carbide forming elements is 0.05 wt % or more and 0.50 wt %
or less.
6. The alloy according to claim 2, characterized in that an amount
of carbide forming elements which have not been carbidized is 0.50
wt % or less.
7. The alloy according to claim 2, characterized in that a compound
phase of the nickel and the carbide forming elements is formed.
8. The alloy according to claim 2, characterized in that the
carbide forming elements are titanium or niobium.
9. A lens holding member comprising the alloy according to claim
1.
10. An optical apparatus comprising the lens holding member
according to claim 9.
11. A method for producing an alloy comprising iron, nickel, and
cobalt which are basic components of a Super Invar alloy,
characterized by comprising: adding carbide forming elements to the
basic components and melt casting the resultant mixture; hot
forging at a predetermined temperature; and precipitating carbides
formed from carbons contained in the alloy and the carbide forming
elements in a base phase by performing a first heat treatment at a
first temperature which is lower than the predetermined
temperature.
12. The method for producing an alloy according to claim 11,
characterized in that the first temperature is 825.degree. C. or
more and 950.degree. C. or less.
13. The method for producing an alloy according to claim 11,
characterized in that a content of the added carbide forming
elements is 0.05 wt % or more and 0.50 wt % or less.
14. The method for producing an alloy according to claim 11,
characterized by further comprising forming a compound phase of the
nickel and the carbide forming elements by performing a second heat
treatment at a second temperature which is lower than the first
temperature.
15. The method for producing an alloy according to claim 11,
characterized by performing a third heat treatment at a third
temperature which is lower than a Curie temperature of the
alloy.
16. The method for producing an alloy according to claim 15,
characterized in that the third temperature is 80.degree. C. or
less.
Description
TECHNICAL FIELD
[0001] The present invention relates to an alloy with a low
coefficient of thermal expansion used in a structural part of a
precision device and the like, and a method for producing the
alloy.
BACKGROUND ART
[0002] Conventionally, there have been publications reporting the
temporal deformation of an Invar alloy (Fe and 36% Ni), which is a
kind of a low thermal expansion metal material (refer to Physics
and Applications of Invar
[0003] Alloys, Honda Memorial Series on Material Science No. 3,
AGING AND PRECIPITATION 1978, Maruzen). This publication uses the
term ".gamma. expansion", and conjectures that a cause for .gamma.
expansion is the carbon which is contained in the alloy.
[0004] Further, there are also examples which point out the
problems of temporal dimensional change over a long period of time
in the metal material used in a structural part of a precision
device (refer to Japanese Patent Application Laid-Open No.
H08-269613). The cause in this case is the result of a release
process of the inner residual stress conferred during production
processes such as a heat treatment and the like. This publication
reports, although it is for a cast iron type low thermal expansion
alloy with carbon content of 0.3 to 2.5 wt %, the low thermal
expansion could be achieved by reducing Ni localization. Concerning
the release of the inner residual stress of the latter publication,
this is a known phenomenon in the production of parts made from
metal, and until now procedures and treatments to suppress this
phenomenon have been carried out in the production processes of
precision parts.
[0005] On the other hand, regarding the former publication, this is
still at the academic stage. Even among persons skilled in the art,
the reality is that this phenomenon is not well known. The reason
for that is considered to be that the absolute amount of the
temporal deformation produced by the phenomenon of .gamma.
expansion is smaller than the temporal deformation produced by the
other common phenomena.
[0006] Further, in recent precision apparatuses which are becoming
substantially more precise, increasingly Super
[0007] Invar alloys, not Invar alloys, are used which reliably have
a coefficient of thermal expansion at temperatures close to room
temperature of less than 1 ppm (=1.times.10.sup.-6)/degree. Of
course, while some inorganic materials have a coefficient of
thermal expansion of less than 1 ppm/degree, the production
processes of such materials are very difficult, such as cutting
processing being impossible and the like. Further, since such
materials are not very tough, the material can be damaged during
the production processes. Moreover, since such materials have a
small thermal conductivity, when a localized temperature
distribution is produced in the member, partial expansion can
occur, which prevents the characteristic of having a small
coefficient of thermal expansion from being fully exploited.
Therefore, there is a need to skillfully use Super Invar alloys
while exploiting their characteristics.
DISCLOSURE OF THE INVENTION
[0008] However, concerning Super Invar alloys, whether a .gamma.
expansion phenomenon occurs in the same manner as in Invar alloys
has not been grasped. Further, for high-precision optical
apparatuses, the performance of the apparatus can gradually
deteriorate over a long period of time due to a change in the
optical path length as a result of temporal deformation of an
important structural part. Therefore, in the investigation process
leading to the present invention, to suppress the respective
well-known phenomena which are causes of the temporal deformation
of a metal material, Super Invar alloys were prepared in which the
usual treatments were diligently performed. Despite this, temporal
deformation still remained, and the amount of this deformation
(temporal deformation amount of 5 ppm annually) was such that it
could not be ignored as a cause for performance deterioration of
optical apparatuses which will continue to become more precise in
the future. If the microdeformation evaluation system which greatly
contributed to the present invention is used, minute amounts of
temporal deformation which have conventionally been overlooked
become clear.
[0009] It is an object of the present invention to provide an alloy
which can suppress as much as possible minute temporal deformation
of a Super Invar alloy, and a method for producing such an
alloy.
[0010] In view of the above-described problems, the alloy of the
present invention includes iron, nickel, and cobalt, which are the
basic components of a Super Invar alloy, and is characterized in
that an amount of a fraction which has not been carbidized in
carbon contained in the alloy is 0.010 wt % or less.
[0011] Further, a method for producing an alloy of the present
invention is a method for producing an alloy including iron,
nickel, and cobalt which are basic components of a Super Invar
alloy, characterized by including:
[0012] adding carbide forming elements to the basic components and
melt casting the resultant mixture;
[0013] hot forging at a predetermined temperature; and
[0014] precipitating into a base phase carbides formed from carbons
contained in the alloy and the carbide forming elements by
performing a first heat treatment at a first temperature which is
lower than the predetermined temperature.
[0015] The present inventors discovered that minute temporal
deformation of a Super Invar alloy is caused by a fraction which
has not been carbidized in carbon contained in the alloy.
[0016] Further, with the alloy of the present invention, minute
temporal deformation of a Super Invar alloy can be suppressed as
much as possible (specifically, 2 ppm (2.times.10.sup.-6) or less
calculated on an annual basis).
[0017] In addition, according to the method for producing an alloy
of the present invention, even if only a trace amount of carbide
forming elements is added, the carbide forming elements effectively
combine with the carbon to form carbides. As a result, the amount
of a fraction which has not been carbidized is carbon contained in
the alloy can be made to be 0.010 wt % or less. Therefore, an alloy
which can suppress as much as possible temporal deformation while
maintaining a low coefficient of thermal expansion can be
produced.
[0018] Further features of the present invention will become
apparent from the following description of exemplary embodiments
with reference to the attached drawings.
BRIEF DESCRIPTION OF THE DRAWINGS
[0019] FIG. 1A is a graph illustrating a stress-strain curve
obtained as a result of a three point bending test.
[0020] FIG. 1B illustrates a schematic diagram of the three point
bending test of FIG. 1A.
[0021] FIG. 2 is a graph illustrating the maximum surface stress of
various composition materials and a permanent strain amount
remaining at that time.
[0022] FIG. 3 illustrates an actually-measured amount of temporal
deformation of the various composition materials.
[0023] FIG. 4 is a schematic diagram illustrating a measurement
system used in the measurement of the amount of temporal
deformation.
[0024] FIGS. 5A, 5B and 5C are photographs illustrating the metal
structure of a high strength material 8-2 and a graph illustrating
the results of dissolution extraction analysis.
[0025] FIG. 6 is a graph illustrating a relationship between the
amount of free carbon of the various composition materials
determined from the results of the dissolution extraction analysis
and the temporal deformation amount.
[0026] FIG. 7 is a graph illustrating the proportion of fixed
carbon content (wt %) with respect to the total carbon content (wt
%), calculated using the fixed carbon content (wt %) determined
from the results of the dissolution extraction analysis of FIG.
5C.
[0027] FIG. 8 is a graph illustrating a relationship between the
heat treatment temperature and the amount of free carbon.
[0028] FIG. 9 is a graph illustrating the coefficient of thermal
expansion of the respective Super Invar alloy composition materials
used in the present investigation.
[0029] FIG. 10 is a graph illustrating the results of measuring the
temperature and displacement (dimensional change of a test piece)
of a high strength material (C: 0.118%) by a thermal expansion
meter.
[0030] FIG. 11 is a cross sectional view of a lens barrel using a
frame member 8 of the alloy of the present invention.
BEST MODE FOR CARRYING OUT THE INVENTION
[0031] The present invention will now be described in more detail
with reference to the drawings. Further, the same structural
elements are in principal provided with the same reference numeral,
and thus a description thereof is omitted.
(Alloy)
[0032] The alloy of the present invention includes iron, nickel,
and cobalt, which are the basic components of a Super Invar alloy.
This alloy is characterized in that the amount of a fraction which
has not been carbidized in carbon unavoidably included in the alloy
is 0.010 wt % or less.
[0033] Super Invar alloys have iron, nickel, and cobalt as basic
components, with a coefficient of thermal expansion which is even
lower than that of Invar alloys (Fe 63.5 wt %, Ni 36.5 wt %). The
basic composition of a Super Invar alloy is Fe 63.5 wt %, Ni 31.5
wt %, and Co 5.0 wt %, obtained by replacing 5 wt % of the Ni of
the Invar alloy with Co.
[0034] In the present specification, "basic components" means the
essential components of a Super Invar alloy.
[0035] In the alloy of the present invention, the content ranges of
the above-described basic components may be illustrated as
follows.
[0036] The Co content may be 2.0 wt % or more and 8.0 wt % or less,
and can further be 3.0 wt % or more and 7.0 wt % or less.
[0037] The Ni content may be 30.0 wt % or more and 38.0 wt % or
less. Here, when the below-described carbide forming elements, such
as Ti and Nb are added in an amount higher than the minimum
necessary amount to fix the carbons, excess carbide forming
elements form a compound with the Ni. Thus, in such a case, a
larger amount of Ni has to be included, taking into consideration
the amount of Ni forming a compound with the excess carbide forming
elements. Therefore, a desirable upper limit is 38.0 wt %. However,
when the minimum necessary amount of the carbide forming elements
is added, the upper limit is about 34.0 wt %.
[0038] The Fe content may be 62.0 wt % or more and 68.0 wt % or
less, and can further be 63.0 wt % or more and 67.0 wt % or
less.
[0039] The present inventors prepared a Super Invar alloy which had
been subjected to common treatments for suppressing the respective
well-known phenomena which are causes of temporal deformation of
metal materials such as an Invar alloys and the like. Despite this,
there was still temporal deformation of about 5 ppm annually. While
this will be described in more detail below, the present inventors
discovered that the cause of this temporal deformation is the
amount of free carbon. The above-described common treatments,
although will be described in more detail below, include e.g. slow
cooling (furnace cooling) and a sub-zero treatment.
[0040] Here, in the present specification, non-carbide-forming
non-attached carbon of carbon contained in the alloy is called
"free carbon". Specifically, a "free carbon" means a solid solution
carbon such as an interstitial carbon. It could be conjectured from
the results of the present experiment that free carbon is a cause
of temporal deformation due to the free carbon intruding into the
lattice and moving in the lattice.
[0041] If the amount of free carbon is 0.010 wt % or less, the
temporal deformation could be suppressed to 2 ppm or less
calculated on an annual basis. Further, the amount of free carbon
is more desirably 0.005 wt % or less. A lower limit of 0.0005 wt %
or less could be realized.
[0042] In the alloy of the present invention, the total carbon
content in the alloy, specifically, the carbon content which is
unavoidably contained in the alloy, may be 0.010 wt % or less. When
a Super Invar alloy is obtained by general industrial processes,
the total carbon content is often about 0.010 wt % to 0.030 wt % or
less. Thus, it was learned that even if the total carbon content is
0.010 wt % or more, as long as the amount of free carbon is 0.010
wt % or less, temporal deformation can be effectively
suppressed.
[0043] To ensure that as few of the carbons which are unavoidably
present in the alloy as possible are not free carbon, the alloy of
the present invention may include carbide forming elements. In the
present specification, "carbide forming elements" are an element
which is capable of combining with carbon in the alloy to form a
carbide. At least a part of the carbide forming elements combines
with carbon in the alloy to form a carbide, and the formed carbide
is dispersed in the base phase. As long as a carbon atom is fixed
as a carbide, the amount of free carbon can be decreased.
[0044] It is desirable to include a larger amount of carbide
forming elements than the total carbon content based on an atomic
percent comparison, otherwise free carbon which is not fixed as
carbides will remain.
[0045] The carbide forming element content is desirably 0.05 wt %
or more. This is because such a content allows an amount of free
carbon equivalent to 0.01 wt % of the total carbon content to be
fixed. On the other hand, to maintain the coefficient of thermal
expansion to less than 1 ppm/.degree. C., the carbide forming
element content is desirably 0.50 wt % or less. The specific
significance of these values will be described below. Further, the
lower limit is more desirably 0.10 wt %. The upper limit is more
desirably 0.30 wt %. The reason for this is as follows. If the
carbide forming element content is too low, the carbons in the
alloy cannot all be fixed as carbides, so that the temporal
deformation cannot be effectively suppressed. On the other hand, if
the carbide forming element content is too high, some of the
carbide forming elements do not combine with the carbons and are
left over, which increases the coefficient of thermal expansion. As
a result, the desirable characteristics of the Super Invar alloy
cannot be exploited.
[0046] The amount of carbide forming elements which do not form a
carbide is desirably as low as possible. If this amount is 0.50 wt
% or less, the low coefficient of thermal expansion which is
characteristic of a Super Invar alloy can be maintained.
[0047] In the present invention, it is desirable that a compound
phase is formed between the nickel and the carbide forming
elements. Carbide forming elements which, as described above, do
not form carbides and are left over, combine with the Ni, which is
a basic component of the alloy. As a result, an increase in the
coefficient of thermal expansion can be suppressed, and the
strength of the alloy can be increased.
[0048] Examples of the carbide forming elements include titanium
(Ti) and niobium (Nb).
(Lens Holding Member, Optical Apparatus)
[0049] The alloy of the present invention can suitably be used for
a lens holding member. FIG. 11 is a cross sectional view of a lens
barrel using a frame member 8 of the alloy of the present
invention. In FIG. 11, a lens 7 made of low thermal expansion (0.6
ppm/degree) quartz is fixed on the frame member 8 having the same
coefficient of thermal expansion. Further, an outer casing 9
supports the frame member.
[0050] Since the coefficient of thermal expansion is the same as
the lens, light passing through the lens (the optical path) does
not change even if a change in temperature occurs.
[0051] Further, since the frame member uses the alloy of the
present invention, it hardly undergoes temporal deformation, so
that the frame member does not impart deformation to the lens which
is fixed thereto. Therefore, since optical errors, such as
aberration, do not occur over a long period of time, the
characteristics of the apparatus using this barrel also do not
change.
[0052] Desirable examples of optical apparatuses for which the
advantageous effects of the present invention would be appreciated,
and which require nanometer level precision include an exposure
apparatus used in manufacturing semiconductor devices, but are not
limited thereto, may also used for e.g. an optical apparatus which
is used in space.
(Alloy Production Method)
[0053] The alloy of the present invention can be produced as
follows, for example.
[0054] If the total carbon content can be suppressed to 0.010 wt %
or less, obviously the amount of free carbon will also be in this
range. Although suppressing the total carbon content to 0.010 wt %
or less can be achieved by exploiting sophisticated refining
techniques such as VAR (vacuum arc remelting) and ESR (electroslag
remelting), this is difficult with ordinary industrial process.
[0055] Further, even for a Super Invar alloy having a total carbon
content of 0.010 wt % or more, by adding carbide forming elements
such as Ti, Nb, and the like, melt casting the resultant mixture,
and hot forging at a predetermined temperature, the carbons can be
fixed by the Ti and the like, which allows an alloy with suppressed
free carbon to be obtained.
[0056] However, as described above, it is desirable to avoid adding
the carbide forming elements in an excessive amount, and in an
amount as small as possible. Further, while it is desirable for all
of the carbide forming elements to react with the carbons in just
the right proportion, this is difficult in actual practice as the
amount of carbons which is unavoidably present varies. Thus, it was
learned that to obtain the carbides from the added carbide forming
elements as efficiently as possible, the following production
method may be carried out.
[0057] First, the carbide forming elements are added to the basic
components, and the resultant mixture is melt cast by a 50 kg
vacuum induction melting furnace.
[0058] Then, the mixture is hot forged at a predetermined
temperature. The predetermined temperature may be 1000.degree. C.
or more and 1100.degree. C. or less.
[0059] Further, by carrying out a first heat treatment at a first
temperature which is lower than this predetermined temperature, the
carbides formed from the carbons contained in the alloy and the
carbide forming elements can be precipitated in the base phase.
[0060] The first temperature does not have to be a constant
temperature, and the first heat treatment may be carried out by
holding a temperature which gradually decreases. However, it is
desirable to hold the temperature in a predetermined range. This is
because if the temperature is too high, the formed carbide
decomposes, while if the temperature is too low, it is difficult
for the carbide forming elements and the carbons to bond. From such
a perspective, the first temperature is desirably 825.degree. C. or
more and 950.degree. C. or less.
[0061] The significance of this production method of the present
invention is that the method does not reduce the carbon content by
utilizing an expensive refining technique, but rather suppresses
the undesirable effects of the carbon atoms without an increase in
the coefficient of thermal expansion even if a small amount of
carbon atoms is unavoidably present during a typical melt and
casting production process.
[0062] Further, after the first heat treatment, the compound phase
of the nickel and the carbide forming elements may be formed by
carrying out a second heat treatment at a second temperature which
is lower than the first temperature. As a result, excess carbide
forming elements which did not contribute to carbide formation can
be suppressed as much as possible from becoming a cause of an
increase in the coefficient of thermal expansion, thereby allowing
the strength of the alloy to be improved.
[0063] The second temperature may be 700.degree. C. or more and
750.degree. C. or less.
[0064] Further, after the second heat treatment, a third heat
treatment may be carried out at a third temperature which is equal
to or higher than room temperature to lower than the Curie
temperature of the alloy. As a result, free carbon which could not
turn into a carbide can diffuse to a stable position in the Super
Invar alloy. By applying a so-called artificial seasoning effect,
the temporal deformation can be further reduced.
[0065] The third temperature may be 25.degree. C. or more and
150.degree. C. or less. This upper limit may further be 120.degree.
C.
EXAMPLE
(Concerning the Causes of Temporal Deformation)
[0066] First, using members of Super Invar alloys and Super Invar
alloy equivalent materials, a procedure from which the causes of
temporal deformation, which conventionally has been a problem,
could be recursively inferred was constructed. Specifically, unlike
the causes of typical temporal deformation, a hypothesis that
carbon atoms have some sort of effect on temporal deformation was
constructed.
[0067] The following three procedures were carried out to decrease
temporal deformation: (i) decrease the amount of carbon atoms which
are a cause to a level which is not a problem; (ii) prevent the
unavoidably contained carbon atoms from diffusing by fixing the
carbon atoms by some sort of procedure; and (iii) carry out a metal
structure stabilizing treatment (so-called "seasoning") for a
practical period of time.
[0068] As a result of this investigation, it was discovered that
all of these procedures were effective in reducing temporal
deformation. This result is related to the fact that the carbon
atoms are a cause. Further, this result indicates that containing
carbon atoms in itself is not a problem, but that the problem is
what state the carbon atoms are present in the metal structure. In
other words, it is better not to allow the carbon atoms to diffuse
in the lattice at temperatures close to room temperature.
[0069] The low thermal expansion alloy and production method
thereof of the present embodiment employ what is currently thought
to be the best measures. Specifically, in the melt casting step,
care is taken to prevent the incorporation of impurities. The
content of unavoidable carbons is suppressed, and a trace amount of
an element is added for forming a carbide by strongly bonding with
the carbon atoms. Further, a suitable heat treatment is carried out
to effectively precipitate the carbide, so that the amount of free
carbon (meaning solid solution carbons or interstitial carbons) is
reduced.
[0070] The various composition materials formed from adding the
carbide forming elements to the original Super Invar alloy are
herein referred to as a "Super Invar alloy equivalent
material".
(Investigation of the Various Phenomena which are Causes of
Temporal Deformation)
[0071] Super Invar alloys which are used for precision optical
apparatus parts have a basic component elemental composition of
31.5% Ni (nickel), 5% Co (cobalt), and a balance of Fe (iron).
[0072] If the dimensional change of the material is measured at a
constant temperature (23.+-.0.01.degree. C.) controlled with high
precision, it was discovered that at an initial stage after the
part was produced, extension of about 0.5 ppm calculated as a
monthly basis continued. However, this is uneven depending on the
production lot of the alloy, and in some lots there was greater
extension. Having an assumption that the cause was due to
thermally-activated process phenomena, such as diffusion and the
like, the rate of change should decrease with the passage of time.
Even still, it was learned from the results of measuring the
dimensional change for several alloys for several tens of days, the
deformation over about one month from the point where the final
treatment was finished may be taken as linear change versus time
elapsed, and that the extension calculated as a month could be
estimated even by measurements of about 5 to 10 days.
[0073] However, it is often difficult to specify the cause of
dimensional change in a metal structure based on metal structural
changes. If improvements could be found by a symptomatic treatment
method, whatever responded to that treatment would normally be
assumed to be the cause. Even currently, where high-resolution
electron microscopes are developed, micro-level changes at the
atomic level can be said to be impossible to directly detect in
their original state even for metal structural changes. Therefore,
the cause must be conjectured based on a comprehensive
consideration of information and the like from other checks,
experiments, and analysis. In the present invention as well, the
following respective phenomena were first investigated as the cause
of this temporal deformation, and then the cause was narrowed down
by a process of elimination.
(1--Concerning Reduction of Spontaneous Magnetization)
[0074] In Super Invar alloys, the same as Invar alloys, the normal
atomic distance determined by the thermal vibration of the atoms is
enlarged due to spontaneous magnetization specific to these alloys.
As the temperature increases, the level of this spontaneous
magnetization is reduced. Therefore, this spontaneous magnetization
acts to cancel out the normal enlarged amount of atomic distance
caused by an increase in temperature. As a result, Super Invar
alloys exhibit a very small coefficient of thermal expansion at the
Curie temperature or less, and at more than the Curie temperature
the coefficient of thermal expansion returns to normal. Therefore,
if spontaneous magnetization does decrease over time, the volume
will move in a decreasing direction, which is the reverse of the
presently-occurring extension (expansion) phenomenon. Thus, this
phenomenon was not considered.
(2--Concerning Precipitation Phenomenon)
[0075] At close to room temperature, all of the component elements
of a Super Invar alloy, Fe, Ni, and Co have a very small
self-diffusion coefficient, and thus the diffusion of these atoms
would be difficult to consider.
[0076] However, while atoms such as carbon which have a small
atomic radius do exhibit interstitial diffusion even at room
temperature, from the results of diffusion, it cannot be confirmed
whether some kind of stable phase (e.g., graphite, cementite and
the like) is generated.
[0077] Further, because this is an important point it will be
described in more detail below, but concerning temporal dimensional
change, in the range of from room temperature, which is the Curie
temperature or less, to 210.degree. C., reversible expansion, which
was not normal thermal expansion, and contraction were found. It
cannot be denied that there is a possibility of precipitation even
at such a low temperature range. However, that precipitated
material is unconceivable to re-dissolve and again form a solid
solution. Therefore, since this phenomenon is not reversible at the
above-described temperature range, it was excluded as the cause of
temporal deformation.
(3--Concerning Release of Inner Residual Stress)
[0078] This phenomenon is a process which has inner stress as a
driving force, in which as a result of the movement and
disappearance of dislocations, which are a kind of lattice defect,
dislocation density decreases. Thus, volume can be thought to move
in a decreasing direction. However, just in case, a temporal
deformation comparison measurement was carried out, using the same
materials, between a material obtained by air cooling from
315.degree. C. and a material obtained by furnace cooling, yet no
difference was found. Therefore, this phenomenon was also
excluded.
(4--Concerning Martensitic Transformation)
[0079] This phenomenon involves an increase in volume. However, the
unique, high-precision dimensional change measurement performed in
the present invention was carried out in a location which was
controlled at .+-.0.01.degree. C. of 23.degree. C. However, for the
present alloy, it can be considered that this transformation will
not occur at a constant temperature. This is because, unlike
typical tool steel and the like, the stable phase at room
temperature for the present nickel-containing, iron-based low
thermal coefficient alloy is originally an austenite phase.
Further, there is no change in stress applied on the alloy during
the dimensional change measurement. Therefore, stress-induced
transformation also does not progress. However, despite this, using
the same materials, a material which was subjected to a treatment
at a low temperature (sub-zero treatment) of -10.degree. C. and a
material which was not subjected to such a treatment were compared.
From the results, no difference was found in the temporal
deformation amount. Therefore, this phenomenon was also
excluded.
(5--Concerning Creep)
[0080] Considering the amount of temporal deformation which is a
problem, creep deformation caused by the sample's own weight may
also contribute. For typical steel, which generally is a
body-centered cubic lattice, the diffusion of interstitial elements
tends to occur even at room temperature, so that strain aging
dynamically progresses, which makes it difficult for creep
deformation to occur. Concerning this point, for this alloy which
is a face-centered cubic lattice, the diffusion rate of carbon is
smaller than that for steel, so that it is said that strain aging
at temperatures close to room temperature does not progress easily,
and creep occurs more easily.
[0081] Accordingly, the strength in the microstrain range (at a
permanent strain of 2 ppm=about 0.0002%) of the various alloys was
measured. It can be said that dislocations are less likely to
diffuse and creep resistance is larger as the strength of the alloy
at this strain amount is larger. Among the samples which were
measured for temporal deformation, as expected the high-carbon
material (C: 0.118 wt %) had a larger strength than the low-carbon
material (C: 0.002 wt %). Thus, if creep is the cause, high-carbon
materials should have a smaller amount of temporal deformation.
Here, in the measured samples, other than carbon, the contained
components were exactly the same. Various samples were also
prepared by melt production from both electrolytic iron and
electrolytic nickel, and then further adding carbon for the
high-carbon materials. The results of the temporal deformation
measurement will be described below.
(6--Concerning .gamma. Expansion)
[0082] The gaps in an austenite phase (.gamma. phase) face-centered
cubic (fcc) lattice are the centers of an octahedron and the
centers of a tetrahedron, the latter one having geometrically even
smaller gaps. For example, if carbon atoms which were in gaps in
the center position of an octahedron diffuse into gaps in the
center position of a tetrahedron the volume may expand. Due to the
existence of spontaneous magnetization, carbon atoms can stably be
present even in positions where the geometric gap is small. If the
mechanism of .gamma. expansion is based on diffusion between nearby
gaps among the lattices of carbon atoms such as described above,
then the temporal deformation amount can be expected to decrease
for a low-carbon material having few free carbon. Further, since
the deformation is not produced by production of a compound or the
like, the mechanism can be expected to be reversible with respect
to the temperature change of the temperature range mentioned above.
While this will be described in more detail below, based on a
series of investigations for the present invention, the cause of
temporal deformation could be specified as being this
phenomenon.
(Alloys According to the Examples)
[0083] Table 1 is a list of various Super Invar alloys and Super
Invar alloy equivalent materials which were prepared for measuring
the temporal deformation amount.
TABLE-US-00001 TABLE 1 Prediction of Separate Various Component (wt
%) Causes of Temporal Composition Total Other Main Deformation
Materials Carbon Compo- Creep .gamma. Expansion Name Code Content
nent (s) Hardness Ease Ease Comparative Conventional STD 0.012%
About 120 HV This This Example 1 Material conventional conventional
material is material is the standard the standard Example 1
Low-Carbon LC 0.002% About 120 HV Easy for Difficult Material creep
to for .gamma. occur expansion to occur Comparative High-Carbon HC
0.118% About 120 HV Difficult Easy for .gamma. Example 2 Material
for creep to expansion to occur occur Example 2 High 8-1 0.013% Ti
3.0% Equivalent Very Difficult Strength DA to 372 HV difficult for
.gamma. Material for creep to expansion to occur occur Example 3
High 8-2 0.013% Ti 3.0% Equivalent Very Difficult Strength STA to
389 HV difficult for .gamma. Material for creep to expansion to
occur occur Example 4 High 9-1 0.012% Nb 3.9%, Ti Equivalent Very
Difficult Strength DA 0.6% to 287 HV difficult for .gamma. Material
for creep to expansion to occur occur Example 5 High 9-2 0.012% Nb
3.9%, Ti Equivalent Very Difficult Strength STA 0.6% to 252 HV
difficult for .gamma. Material for creep to expansion to occur
occur Example 6 Nb-Added N 0.017% Nb 0.24% Difficult Difficult
Material for creep to for .gamma. occur expansion to occur Example
7 Ti-Trace T 0.023% Ti 0.08% Difficult Difficult Amount Added for
creep to for .gamma. Material occur expansion to occur
[0084] Table 1 simultaneously illustrates, in a case that the cause
of temporal deformation was creep, whether it is predicted that the
various materials are not as susceptible to temporal deformation as
the conventional material, or it is predicted that the various
materials are more susceptible to temporal deformation than the
conventional material. Table 1 also similarly illustrates these
cases for when .gamma. expansion was the cause.
[0085] Although the contained carbon content for Super Invar alloy
equivalent materials thoroughly containing Ti and Nb (high strength
materials, Examples 2 to 5) was basically the same as for the
conventional material, the carbons would be expected to be carbides
such as TiC and NbC would be expected (the analysis results will be
described below). Thus, since the number of carbon atoms which can
actually move should be smaller, the temporal deformation would be
expected to decrease in either case that the cause is creep or
.gamma. expansion.
[0086] Hardness is written as an average value, because a hardness
of "about 120 HV" depends on the location of the hardness test.
Further, "equivalent to 372 HV" is actually a value obtained by
converting the value of the results found by testing with a
Rockwell hardness testing machine to a Vickers hardness.
[0087] In addition, the DA material of the various composition
materials means a material obtained by hot forging, then air
cooling, and then performing an aging treatment. Further, an STA
material means a material obtained by hot forging, then performing
a solution treatment, then water cooling, and then performing an
aging treatment.
(Method for Producing the Alloys According to the Example)
[0088] The conventional material, low-carbon material, high-carbon
material, and Nb-added material (Examples 1, 6, Comparative
Examples 1 and 2) were produced by undergoing melt casting by a
vacuum induction melting furnace, air cooling after hot forging at
1000.degree. C., water cooling after a solution treatment by
holding at 830.degree. C. for 2 hours, performing a stress release
treatment by holding at 315.degree. C. for 3 hours, and performing
a stabilizing treatment at 98.degree. C. for 48 hours (so-called
artificial seasoning, the third heat treatment).
[0089] Further, the high-strength materials (Examples 2 to 5) were
produced by undergoing melt casting by a 50 kg vacuum induction
melting furnace, hot forging at 1000.degree. C., performing an
aging treatment at 720.degree. C. for 6 hours (second heat
treatment), then performing a stress release treatment at
315.degree., and performing a stabilizing treatment at 98.degree.
C. for 48 hours.
[0090] Further, the Ti-trace added material (Example 7) was
produced by undergoing hot forging at 1000.degree. C., holding at
900.degree. C. for 2 hours (first heat treatment), then gradually
cooling to 830.degree. C., air cooling from that temperature to
room temperature, then performing an aging treatment at 720.degree.
C. for 6 hours (second heat treatment), then performing a stress
release treatment at 315.degree., and performing a stabilizing
treatment at 98.degree. C. for 48 hours (third heat treatment).
(Contained Components of the Alloys According to the Example)
[0091] Table 2 lists the contained components of the various
composition materials.
TABLE-US-00002 TABLE 2 Provided Test Materials in the Present
Invention Weight (%) C Si Mn P S Cu Ni Cr Comparative Conventional
0.012 0.24 0.33 0.003 0.004 <0.01 32.69 -- Example 1 Material
Example 1 Low-Carbon 0.002 0.25 0.34 <0.002 0.001 <0.01 32.8
<0.01 Material Comparative High-Carbon 0.118 0.25 0.35 <0.002
0.001 <0.01 32.8 <0.01 Example 2 Material Example 2 High
Strength 0.013 0.26 0.36 <0.01 <0.01 <0.01 36.9 0.01 DA
Material 8-1 Example 3 High Strength Material 8-2 Example 4 High
Strength 0.012 0.25 0.35 <0.01 <0.01 <0.01 36.9 <0.01
DA Material 9-1 Example 5 High Strength Material 9-2 Example 6
Nb-Added 0.017 0.25 0.25 0.005 0.012 <0.01 33.76 0.02 Material
Example 7 Ti-Trace 0.023 0.25 0.35 <0.002 0.001 <0.01 32.59
<0.01 Amount Added Material Weight (%) Mo Co Ti Nb O N Fe
Comparative Conventional -- 4.86 -- -- Balance Example 1 Material
Example 1 Low-Carbon <0.01 4.84 -- -- Balance Material
Comparative High-Carbon <0.01 4.92 -- -- Balance Example 2
Material Example 2 High Strength -- 4.88 3.0 -- Balance DA Material
8-1 Example 3 High Strength Material 8-2 Example 4 High Strength --
4.93 0.6 3.9 Balance DA Material 9-1 Example 5 High Strength
Material 9-2 Example 6 Nb-Added -- 3.75 <0.01 0.24 Balance
Material Example 7 Ti-Trace -- 5.10 0.08 -- 0.0008 0.0006 Balance
Amount Added Material
(Conventional Material, Low-Carbon Materials, High-Carbon
Materials)
[0092] The low-carbon materials and high-carbon materials were each
produced using high-purity materials produced by electrolysis as
the raw materials by changing only their carbon contents, and
having identical amounts for the other elements.
(Concerning the High-Strength Materials)
[0093] Each of the high-strength materials had the same carbon
content as the conventional material, and was made to thoroughly
contain Ti and/or Nb.
[0094] The results of the hardness testing were that after the hot
forging, materials 8-1 and 9-1 respectively had a hardness of 36.6
HRC (equivalent to 360 HV) and 25.7 HRC (equivalent to 270 HV), and
that after the solution treatment, materials 8-2 and 9-2
respectively had a hardness of 33.3 HRC (equivalent to 330 HV) and
18.9 HRC (equivalent to 234 HV). Since the hardness illustrated in
Table 1 is a value after the aging treatment (720.degree.
C..times.6 hours), all of the hardness values were increased by
carrying out the aging treatment for precipitation phase
formation.
[0095] Further, Ti addition had a larger effect on increasing
hardness than Nb addition. Ni forms a precipitation phase, which is
a compound phase with the Ti and Nb, by the aging treatment.
Therefore, the Ni content is increased based on the carbide forming
element content so that the matrix portion becomes a composition
which can exhibit an Invar effect (effect in which the coefficient
of thermal expansion decreases) to the maximum extent.
(Concerning the Nb-added Material)
[0096] An Nb-added material is a material obtained by adding a
small amount (0.24%) of Nb to a conventional material. The purpose
of this material is to reduce .gamma. expansion, which is said to
be a cause of the above-described diffusion of carbon, by making
unavoidably-contained carbons bond with the Nb, whereby the carbons
are fixed as a carbide.
[0097] The atomic weights of niobium and carbon is respectively
92.9 and 12. As described below, the niobium carbide confirmed in
the present invention is NbC, and thus the atoms bond at a 1:1
ratio. Therefore, unless 7.74 times or more (=92.9/12) of niobium
than carbon is added, free carbon will remain. Further, since it is
difficult for all of the Nb to bond with the carbon atoms, it is
desirable to add more than 7.74 times of niobium than carbon.
However, the addition of more Nb than necessary amount to form the
carbides should be suppressed as much as possible. This is because
residual Nb which does not form a carbide remains, and if this
remaining niobium forms a solid solution in the matrix, the
coefficient of thermal expansion is increased. The same can be said
for the carbide forming elements other than Nb.
[0098] In the present example, since the carbon content is 0.017 wt
%, it is desirable to add 0.13 wt % or more of Nb. Thus, the Nb
content was adjusted to 0.24 wt %.
[0099] Even in this case, like with the above-described high
strength materials, it is desirable to carry out an aging heat
treatment (second heat treatment) on the niobium formed in a solid
solution. This is because the hardness is increased due to the
formation of a precipitation phase of Nb and Ni. This aging heat
treatment is carried out to increase the strength value at the
microstrain (micro yield point). This is because the present
invention is directed to resolving the problem of a temporal
deformation which is very small. The strength value at the
microstrain is the stress of a remaining permanent strain at a
level of 0.0001% (1 ppm). If this micro yield point can be
increased, creep resistance also increases. At the above-described
production conditions, this aging heat treatment is equivalent to
720.degree. C. While the measurements here measured the temporal
deformation in a state where a low thermal expansion alloy was
arranged on a flat surface, in an actual apparatus, a larger stress
than this state may be applied on a member which is formed from a
low thermal expansion alloy. Therefore, it is even more desirable
to improve creep resistance. Further, increasing the hardness of
the alloy by even a little bit achieves desirable results also in
the cutting and grinding processing. This is because it is easier
to remove the swarf and clogging of the grinding stone is
reduced.
[0100] Obviously, not only is the .gamma. expansion of the
respective high strength materials reduced, but an effect in
suppressing creep deformation can also be expected. This is because
the strength value at the microstrain is increasing as a result of
the precipitation hardening and the solid solution hardening by the
carbide forming element.
(Concerning Ti-Trace Amount Added Material)
[0101] This material has the minimum necessary amount of Ti added
thereto in consideration of a case where ideally all of the carbons
unavoidably present in the alloy are used in the formation of
titanium carbides.
[0102] Since Ti in solid solution reduces spontaneous
magnetization, and effectively increases the coefficient of thermal
expansion, the minimum necessary amount of Ti to maintain a low
coefficient of thermal expansion was used.
[0103] However, while as a result the temporal deformation could be
reduced by a considerable amount, the component blend could not be
adjusted to a Ti value which was just enough. Namely, since the
carbon content was 0.023 wt %, even if all of the added 0.08 wt %
of Ti ideally formed into carbides, the Ti was still not quite
enough. Specifically, since the atomic weight of titanium is 47.9,
at least 0.003 wt % (=0.23-0.08.times.12/47.9) of free carbon,
which admittedly is a tiny amount, will nonetheless remain.
[0104] By carrying out a suitable carbide forming heat treatment
(first heat treatment; in the experiment results, 825.degree. C. or
more and 950.degree. C.), a considerably large proportion of 74% of
the contained Ti becomes a carbide at the optimum temperature of
900.degree. C., so that 0.008 wt % of free carbon remained. With
about this amount of free carbon, the actual temporal deformation
could be reduced as described below.
[0105] Further, the coefficient of thermal expansion also shows
almost no difference depending on the cutting direction of the test
piece. This value was 0.42 to 0.47 ppm/degree in the range of 18 to
28.degree. C. This measurement was carried out using the TMA 8310
manufactured by Ulvac-Riko, Inc., at a rate of temperature increase
of 5.degree. C./minute.
[0106] Further, with the normally used dissolving and casting
process, about 0.03 wt % of carbon is unavoidably contained. To
reduce this level to 0.008 wt % is difficult in actual practice. By
adding a trace amount of Ti, the amount of carbon which influences
the temporal deformation could be reduced to 0.008 wt %, while
hardly changing the coefficient of thermal expansion and with the
same level of production costs.
(Specifying the Main Causes of Temporal Deformation)
[0107] FIG. 1A is a graph illustrating a stress-strain curve
obtained as a result of a three point bending test. In FIG. 1A, the
vertical axis represents the center concentrated load (kN), and the
horizontal axis represents the maximum strain (ppm). Further, in
FIG. 1A, the circles represent a measurement point when adding the
load, and the triangles represent a measurement point when removing
the load. FIG. 1B illustrates a schematic diagram of the three
point bending test. Since it is the microstrain portion which is of
consequence in creep deformation, the portions with a large strain
are not illustrated. In FIG. 1A, as an example of the conventional
material, an increase in the permanent strain amount remaining
while the maximum center load is gradually increased is
illustrated.
[0108] As the test piece, the various composition materials
illustrated in Table 1 were prepared in a 30.times.30.times.339 mm
size. The test piece was supported near either end with a span
interval of 250 mm. A concentrated load was applied on a point in
the center section, and the strain amount was determined from the
output of a strain gauge stuck to the center section of the
opposite surface. The testing temperature was room temperature, and
the subsequent temperature when measuring the temporal deformation
was roughly the same. The loading velocity was set to a small level
of 0.5 mm/minute so that the effects of elastic after-effects could
be ignored.
[0109] First, the maximum load was set to a small level, and the
load was repeatedly added and removed. Then, while gradually
increasing the maximum load, the maximum load at which remaining
permanent strain could initially be detected was 12 kN. The
remaining permanent strain (a) at that stage was 3 ppm. However, in
the present test, since the strain gauge is attached to a part of
the test piece where the maximum strain is produced, and the strain
amount is measured at that portion, the average strain amount of
the whole test piece will be an even smaller value.
[0110] Subsequently, the remaining permanent strain (b) when the
maximum load was set to 15 kN, that load was applied, and then the
load was removed, was 10 ppm. In this case, if the maximum load of
15 kN was applied at the start, a permanent strain of 13 ppm should
remain.
[0111] FIG. 2 is a graph illustrating the maximum surface stress of
the various composition materials of Table 1 and the permanent
strain amount remaining at that stage at that location. In FIG. 2,
the maximum load was converted into the maximum surface stress from
the maximum load determined by the three point bending test of FIG.
1A and the permanent strain amount remaining at that stage, and the
permanent strain amount used the cumulative value of the remaining
strain amount produced until that stress.
[0112] Typically, the permanent strain amount used when displaying
a strength value is 0.2% (200 ppm). However, in the present
investigation, the strength value for a very small permanent strain
amount was determined. In high-precision apparatuses, the
deformation amount which influences the performance of the
apparatus is, considering the use period and maintenance period as
the lifecycle of the apparatus, about 10 ppm for the initial year.
Thus, the measurement was carried out having a resolution of 1 ppm,
which is one order smaller.
[0113] From these results, if the cause of temporal deformation is
the phenomenon of creep deformation, the high strength materials
(Examples 2 to 5) should have a markedly smaller temporal
deformation. Further, for the three test pieces on the left of FIG.
2, in which only the carbon content is different (low-carbon
material: Example 1; conventional material: Comparative Example 1;
and high-carbon material: Comparative Example 2), the strength
value increased as the carbon content increased. As a result,
similarly, supposing that temporal deformation is produced by creep
deformation, the amount of temporal deformation should decrease for
high-carbon strength materials.
[0114] FIG. 3 illustrates the actually-measured amount of temporal
deformation of the various composition materials. In FIG. 3, the
amount of temporal deformation per month is illustrated for when
the various nickel-containing, iron-based low thermal expansion
alloys of Table 1 are assumed to undergo linear deformation for 1
month. From the results, stating the conventional
material(Comparative Example 1) as a standard, the high-carbon
material (Comparative Example 2) showed very large amounts of
temporal deformation. Thus, a comparison among materials which have
different carbon contents showed that the amount of temporal
deformation increased for the high-carbon materials. As long as the
high-carbon material was produced in an ordinary manner, the alloy
composition contained a large amount of carbon which would be
impossible as the unavoidable carbon content. To further pursue the
causes of temporal deformation, along with a low-carbon material,
alloys were specially produced which had a very low and a very high
carbon content. In a comparison among the materials with different
carbon contents (Example 1 and Comparative Examples 1 and 2), the
amount of temporal deformation increased for the high-carbon
materials.
[0115] Further, as will be described below, it was discovered that
size hysteresis clearly appeared in the high-carbon materials
during the heating and cooling when measuring the thermal
expansion.
[0116] On the other hand, the respective high strength materials
(Examples 2 to 5) all had a very small temporal deformation. The
high strength materials contained an amount of about 3.0 wt %,
which was considerably larger than the minimum necessary amount to
form the carbides, of carbide forming elements. Therefore, the
strength value should be increased, there should be a marked effect
on reducing the temporal deformation for which creep is a cause,
and there should also be a marked effect on reducing the temporal
deformation for which .gamma. expansion relating to the carbon
atoms is a cause.
[0117] Therefore, while it could not be specified whether the cause
was creep deformation or .gamma. expansion in a comparison of the
results of the temporal deformation for the conventional material
and the high strength materials, considering the results of the
above-described comparison among the materials in which only the
carbon content was changed, the main cause of temporal deformation
can be presumed to be .gamma. expansion.
[0118] The Nb-added materials and the Ti-trace amount added
materials were also considerably improved compared with the
conventional material. Even still, since the amount of free carbon
is slightly more than the high strength materials, the temporal
deformation was slightly larger than that for the high strength
materials. For the conventional material (Comparative Example 1)
and Examples 2 to 7, even though the total carbon content was about
the same, the amount of free carbon for the latter was reduced.
Further, the temporal deformation in Examples 2 to 7 was quite
suppressed. From these results, it was learned that the amount of
the temporal deformation has a close relationship with the amount
of free carbon, and not with the total carbon content.
[0119] Among the common materials used for comparison, quartz
produces hardly any temporal deformation. Further, stainless steel
(SUS 316, 650.degree. C..times.2 hours, then furnace cooling) and
carbon steel (S45C, 800.degree. C..times.2 hours, then furnace
cooling) had a negative value, and thus it can be considered that a
slight inner residual stress remained. It is considered that
.gamma. expansion does not occur in either of these materials.
[0120] FIG. 4 is a schematic diagram illustrating a measurement
system used in the measurement of the amount of temporal
deformation. In this measurement system, a test piece 1 has the
same dimensions as previously used in the three point bending test.
Mirrors 4 and 5 are attached to either end of the test piece 1.
These mirrors 4 and 5 are provided to reflect the laser light from
two laser interferometers (end-measuring machines) 2 and 3. The
laser interferometers 2 and 3 read a dimensional change in the test
piece based on the change in the optical path length, and use the
principles of a so-called "Michelson interferometer".
[0121] Expansion and contraction of the test piece is determined by
the difference in the positional change of either end of the test
piece as measured by the two laser interferometers. The resolution
of this measurement system is 0.2 nm. Since the length of the test
pieces is 339 mm, conversion shows that a resolution of 0.0006 ppm
can be obtained. With such a resolution, the precision is such that
measurement can be sufficiently performed even with a change of
about 1 ppm per year.
[0122] This measurement system is built on a quartz platen 6. The
measurement system is set in an indoor environment with a
temperature controlled within a range of .+-.0.01.degree. C. from
room temperature at 23.degree. C. In actual practice, the
temperature change of the test pieces is even smaller. As a result,
the amount of expansion and contraction due to thermal expansion
can be ignored.
(Method for Measuring the Amount of Free Carbon)
[0123] FIGS. 5A to 5C are photographs illustrating the metal
structure of the high strength material 8-2 as an example and a
graph illustrating the results of dissolution extraction analysis.
By subjecting the surface to a smoothing treatment into a mirror
surface, and then performing a suitable etching treatment to match
the objective, each of the phases can be distinguished. In the
metal structure, other phases different from the base phase are
clearly dispersed. FIG. 5A is a photograph of one of those phases
observed by scanning electron microscopy. A phase of about 5 .mu.m
in size can be confirmed in the center of the photograph. This
phase was subjected to analysis by an Electron Probe Micro Analyser
(EPMA), and titanium and carbon were detected.
[0124] Next, the method for carrying out the dissolution extraction
analysis performed here will be described. First, using 0.5 g of
the alloy material as the anode and platinum as the counter
electrode, electrolysis was carried out at a voltage of 0.2 V in an
electrolysis solution which mixed 10 v/v % of acetylacetone and 1
w/v % of tetramethylammonium chloride. Then, the resultant residue
was subjected to suction filtration in which the residue was passed
through polycarbonate type (PC) filter paper (0.2 .mu.m filter) to
capture the residue. FIG. 5B is a photograph obtained by
observation with scanning electron microscopy in the same manner of
the product captured on this filet paper. When this product was
analyzed by EPMA in the same manner as described above, it was
learned from the results of FIG. 5C that the product was TiC.
[0125] Next, the filter paper was ashed along with the captured
product in a platinum crucible. The resultant product was then
charged with a mixture (solvent) of sodium borate and sodium
carbonate, and the resultant mixture was made to dissolve at
900.degree. C. Then, the mixture was charged with 10 mL of
hydrochloric acid and a few mL of perchloric acid. The dissolved
solution was diluted to a constant volume, and then subjected to
IPC analysis. The proportion of carbides in the dissolved alloy
with respect to the alloy total amount was thus determined, and the
mass of carbon was calculated from the respective atomic weights of
Ti and C. Further, the amount of free carbon was determined by
subtracting the mass of carbon formed as carbides from a total
carbon content analyzed beforehand.
[0126] Here, since the nitrogen in the alloy may also form TiN, a
separate analysis was carried out. However, in this alloy TiN was
not detected. Further, while the high strength materials (9-1 and
9-2, Examples 4 and 5) simultaneously contained Ti and Nb, the
carbon content formed as carbides was determined from the
respective contents and atomic weights. This can be determined in
the same manner even if the carbide forming elements are Ta, Zr and
the like.
[0127] The total carbon content was determined by a combustion
infrared-absorbing method. This analysis is a method in which a
sample is heated to a high temperature in an oxygen stream to
oxidize the contained carbon into carbon dioxide and the like, and
the carbon content is determined by measuring the infrared
absorption.
(Relationship Between the Amount of Free Carbon and Temporal
Deformation)
[0128] FIG. 6 is a graph illustrating the relationship between the
amount of free carbon of the various composition materials
determined from the results of the above-described dissolution
extraction analysis and the temporal deformation amount. In the
graph, the composition materials with the temporal deformation
concentrated near the origin were five composition materials, the
low-carbon material (Example 1) and various high strength materials
(Examples 2 to 5). The other plots also similarly correspond to
Table 1 and FIG. 3.
[0129] At the plots with a low amount of free carbon, there was a
direct relationship between temporal deformation amount and the
amount of free carbon.
[0130] Further, the various high strength materials had a large
portion of their contained carbon fixed as carbides. The proportion
of the amount of free carbon to the total carbon content was 0.0031
wt % or less for all of those materials. Some of the materials had
a proportion which was smaller than that of the low-carbon material
(Example 1, C: 0.002 wt %), which were produced from special raw
materials, such as electrolytic iron, electrolytic nickel or the
like, to have a very low carbon content. As illustrated in FIG. 7,
the amount of free carbon in the high strength materials 8-1
(Example 2), 8-2 (Example 3), 9-1 (Example 4), and 9-2 (Example 5)
was respectively 0.0008 wt % (=0.013-0.0122, hereinafter the same),
0.0005 wt % (=0.013-0.0125), 0.0031 wt % (=0.012-0.0089), and
0.0025 wt % (=0.012-0.0095). Further, similarly, the amount of free
carbon in the Ti-trace added material (Example 6) and Nb-added
material (Example 7) was respectively 0.0080 wt % (=0.023-0.015)
and 0.0069 wt % (=0.017-0.0101).
[0131] These materials also had a very small amount of temporal
deformation. It was learned that if the amount of free carbon is
made to be 0.010 wt % or less, temporal deformation can be
suppressed to a level lower than conventional levels. More
desirably, to suppress the amount of temporal deformation to 1 ppm
or less calculated on an annual basis, the amount of free carbon is
desirably 0.0050 wt % or less.
[0132] Since in the conventional material the carbon is present as
free carbon, unevenness occurs in the amount of temporal
deformation due to unevenness in the unavoidable carbon content
which is included during production. Even for a relatively good
production lot, an expansion of 0.4 ppm calculated on a monthly
basis was shown. In the present invention, temporal deformation can
reliably be reduced to less than that of the conventional material
by causing carbides to form.
[0133] FIG. 7 is a graph illustrating the proportion of fixed
carbon content (wt %) with respect to the total carbon content (wt
%), calculated using the fixed carbon content (wt %) determined
from the results of the dissolution extraction analysis of FIG. 5C.
From the graph, Ti (Examples 7, 2, 3) tends to bond with the carbon
atoms more effectively than Nb (Examples 6, 4, 5), so that only a
small amount of the carbide forming elements needs to be used.
[0134] Since excess carbide forming elements causes the coefficient
of thermal expansion to increase, from this perspective Nb is
better than Ti. Further, in a comparison among composition
materials similarly added with Nb, Examples 4 and 5, which were
subjected an intermediate heat treatment (second heat treatment) at
720.degree. C. for 6 hours, formed the carbides at a slightly
larger proportion than Example 6, which was not subjected to an
intermediate heat treatment. Therefore, to make the amount of
temporal deformation small and to suppress the coefficient of
thermal expansion to a low level, it is desirable to add the
minimum necessary amount of carbide forming elements, and to carry
out a suitable heat treatment in the carbide formation.
[0135] As described above, in the present experiment, in the
[0136] Super Invar alloys, it was learned that the carbide forming
element having the highest carbide forming performance is Ti.
However, other elements commonly mentioned as carbide forming
elements in steel materials, such as Ta, Zr, W, V, Mo and the like,
can also be considered to provide the effects of the present
experiment.
[0137] Further, while the intermediate heat treatment conditions in
the present experiment were 720.degree. C..times.6 hours, these
intermediate heat treatment conditions were carried out on only the
four high strength materials along with the aging treatment
conditions. The primary objective of this intermediate heat
treatment was to harden the materials by causing a compound phase
of the Ni and the carbide forming elements to precipitate.
(Desirable Method for Producing the Alloy)
[0138] To reliably reduce the amount of free carbon, like the four
high strength materials, a large amount of the carbide forming
element may be added. However, if this carbide forming element is
left over, the coefficient of thermal expansion is increased.
Therefore, it is desirable to reduce the amount of free carbon by
effectively combining as small an amount as possible of the carbide
forming elements with the free carbon. To achieve this, it is
desirable to carry out melt casting, hot forging, and then perform
a predetermined heat treatment (first heat treatment).
[0139] FIG. 8 is a graph illustrating a relationship between the
heat treatment temperature and the amount of free carbon. This
graph was prepared in order to determine the temperature where the
carbides tend to form in the first heat treatment. Using the
Ti-trace amount added alloy (Example 7), the amount of free carbon
in a total of six samples which had been melt cast by a vacuum
melting furnace, hot forged at 1000.degree. C., and then subjected
to a heat treatment by holding at a predetermined temperature
(first temperature) for 2 hours, and one sample which had been
subjected to a solution treatment at 1100.degree. C., was measured.
As the first temperature, the measurement was carried out at
intervals of 25.degree. C. in a temperature range of 825.degree. C.
to 950.degree. C.
[0140] From the results of the present experiment, it was learned
that the first temperature at which the carbides tend to form the
most and the amount of free carbon is reduced is 900.degree. C. For
example, a sample which was heat treated at 900.degree. C. had a
total carbon content of 0.023 wt %, of which the fixed carbon
content was 0.0148 wt %. Specifically, 74% of the total carbon
content was formed as carbides.
[0141] Further, since the Ti-trace added material had a Ti content
of 0.08 wt % and carbon content of 0.023 wt %, even if all of the
Ti ideally combined with the carbons, consideration should be given
to the fact that 0.0030 wt % of free carbon will still remain.
[0142] From the present experiment, it was learned that the
carbides form more easily by carrying out the first heat treatment
at a temperature between at least 825.degree. C. or more and
950.degree. C. or less, and more desirably at 875.degree. C. or
more and 925.degree. C. or less.
(Various Measurement Results)
[0143] FIG. 9 is a graph illustrating the coefficient of thermal
expansion of the respective Super Invar alloy composition materials
used in the present investigation. The coefficient of thermal
expansion of the Super Invar alloy composition materials is the
data obtained simultaneously with the results of the dimensional
change measurement carried out by the thermal expansion measurement
device (Laser Thermal Expansion Meter LIX, manufactured by
Ulvac-Riko, Inc.) used to examine the relationship between
temperature and displacement. The test piece has a size of 6 mm in
diameter and 12 mm in length. The heating and cooling rates were
set to 1.degree. C./minute. Since a slight difference occurs
between the temperature of the actual test piece and the
temperature of a thermocouple arranged near to the test piece in
order to measure the atmosphere temperature, the heating and
cooling was carried out in such a slow manner. The measurement
range was between 30.degree. C. and 50.degree. C.
[0144] As a result, as the amount of solid solution carbon
increased, the coefficient of thermal expansion increased, and as
the content of the carbide forming elements increased, the
coefficient of thermal expansion increased.
[0145] Further, no difference was found in the coefficient of
thermal expansion between the DA materials and STA materials. The
STA materials were produced to have a uniform metal structure by
rapidly cooling after carrying out a solution treatment. These STA
materials were produced in order to minimize the Ni segregation in
the metal structure as a Super Invar alloy. This is because if
portions are present in the metal structure which have a low Ni
content or a high Ni content, those portions will be excluded from
the composition which was originally meant to have the lowest
coefficient of thermal expansion, so that as a result the
coefficient of thermal expansion of the whole material is
increased. The influence that the respective elements have on the
coefficient of thermal expansion depends on what state those
elements are present in. While this is the same even for other
carbide forming elements such as Nb, for example, even if Ti is
added, it is desirable for the added Ti to bond with the carbon
atoms which are unavoidably present, and thus be present in the
metal structure as TiC. This is because, strength increases, and
since the TiC itself has a small coefficient of thermal expansion,
coherence with the base phase is small, which means that the action
to increase the coefficient of thermal expansion of the whole
material is small.
[0146] However, if the Ti atoms which do not bond with the carbon
atoms are present by replacing the original lattice points of the
Super Invar alloy, the Invar effects are reduced. In any event, the
strength of spontaneous magnetization decreases due to the carbons
and the carbide forming elements being formed in solid solution. Ti
atoms and the like which do not contribute to the bonding with the
carbon atoms become surplus atoms for a Super Invar alloy which
aims to have a low thermal expansion.
[0147] Accordingly, to achieve a low coefficient of thermal
expansion, the surplus elements may be subjected to a suitable
aging heat treatment to increase strength, so that precipitation
phases such as a .gamma.' (gamma prime) phase and a .gamma.''
(gamma double prime) phase are generated. This is because
simultaneously creep resistance can be improved and cutting
processability can also be improved due to stickiness being
decreased.
[0148] In a Super Invar alloy, unless a special element is added
and a suitable heat treatment is carried out, the carbons which are
unavoidably present will be present as free carbon (solid solution
carbons or interstitial carbons). Those carbon atoms promote
temporal deformation and also increase the coefficient of thermal
expansion.
[0149] Concerning how much the carbide forming elements increase
the coefficient of thermal expansion, in the present investigation,
the increase when 0.24 wt % of Nb was added was 0.25 ppm/.degree.
C., and the increase when 3.9 wt % of Nb was added was 3.0
ppm/.degree. C. For a 36% Ni Invar, the coefficient of thermal
expansion is usually 1 ppm/degree. Since this is too high, Super
Invars were originally selected. To keep a low coefficient of
thermal expansion (less than 1 ppm/degree), which is the greatest
characteristic of Super Invars, the above-described value of 0.50
wt % or less was determined from the approximation curve obtained
during the processes of the present investigation which is
illustrated in expression (2), as a condition for having a
coefficient of thermal expansion of less than 1 ppm/degree.
y=0.76x+0.63 (2)
[0150] Here, y denotes the coefficient of thermal expansion in
units of ppm/degree, and x denotes the Nb content in units of wt
%.
[0151] For Ti, the coefficient of thermal expansion can be
determined from the approximation curve illustrated in expression
(3). Here, x is the same value as for Nb.
y=0.71x+0.60 (3)
[0152] While Ti has a slightly larger ratio which increases the
coefficient of thermal expansion, the value is roughly the
same.
[0153] FIG. 10 is a graph illustrating the results of measuring the
temperature and displacement with the above-described thermal
expansion meter when the high-carbon material (Comparative Example
2, C: 0.118%) was heated and cooled at a rate of 1.degree.
C./minute.
[0154] In the measurement results of the first cycle, a
characteristic was that the slope of the thermal expansion curve of
the test piece decreased at around 150.degree. C. This is thought
to be due to the diffusion of the carbon atoms becoming active at
this temperature. If the temperature is heated to 210.degree. C.
(205.degree. C. in the third and fourth cycles) and then cooled to
room temperature, clearly different trajectories are followed
during the heating and during the cooling. Comparing at the same
temperature close to room temperature, a test piece which has
finished the heating and cooling contracted by 25 ppm (hereinafter,
referred to as "amount due to seasoning effect") from its initial
size. Before this measurement, the test piece had been subjected to
a heat treatment at 98.degree. C. for 48 hours. Subsequently, this
measurement was again carried out after 570 hours had elapsed at
room temperature from the measurement. As a result, among the
carbon atoms formed in solid solution, a certain proportion can be
thought to have moved to a stable position at room temperature
(hereinafter, referred to as "room temperature stable position".
Likely to be a tetrahedral interstitial site in a fcc lattice). If
the heating is carried out to 205.degree. C., spontaneous
magnetization almost completely disappears, so that at this
temperature carbon atoms can be thought to move to another position
(hereinafter, referred to as "high-temperature side stable
position". Likely to be an octahedral interstitial site in a fcc
lattice). Further, it is considered that during cooling the
diffusion of carbons does not keep up, whereby the material returns
to room temperature with most of the high-temperature side stable
positions (a position where the carbon atoms are stable at
205.degree. C., not room temperature) as is.
[0155] Thus, it can be assumed that the above-described 25 ppm
contraction occurred as a result of the change from the volume
(expanded state) of the test piece in a state where a certain
proportion of the carbon atoms were moved to a stable position at
room temperature to a volume (contracted state) in a state where
carbon atoms were moved to an unstable position at room temperature
due to the heating and cooling of the first cycle.
[0156] The temperature of 98.degree. C. appears to have been
determined for the purposes of increasing the rate of interstitial
diffusion of the carbon atoms and finishing the treatment at a time
which is practical for an industrial manufacturing process. It was
probably thought that at this temperature, since the decrease in
spontaneous magnetization is still small, most of the carbon atoms
would be at the same position, which is a stable position at room
temperature. For this point, as described below, the diffusion of
carbons from a room temperature stable position to a
high-temperature side stable position appears to start at
80.degree. C. or more in the present investigation. A temperature
of 80.degree. C. is thus thought more suitable than the
above-described artificial seasoning temperature of 98.degree.
C.
[0157] Although the optimum treatment for temporal deformation is
to leave a member at a use temperature for several tens of years
and then fit the member into a product, this is not practical.
However, in the present investigation, there are also materials for
which measurement of the third cycle was performed after leaving at
room temperature for 1349 hours, without carrying out the
98.degree. C. heat treatment between the second cycle and the third
cycle. In this case, the amount due to the seasoning effect was 19
ppm. In other words, during this time temporal deformation of a 19
ppm expansion occurred. Since 1349 hours is not a practical time
period for an industrial treatment, the below-described temperature
range is better for the stabilizing treatment. When the second
cycle measurement was carried out immediately after the first cycle
was finished, the size of the test piece was not changed even after
undergoing heating and cooling. This can be assumed to be because
at the first cycle heating temperature (210.degree. C.), all of the
carbon atoms had diffused to the high-temperature side stable
position.
[0158] In the second cycle, the reason why the curves do not
exactly match during the heating and the cooling is that the
temperature change of the test piece is slower than the temperature
change of the control thermocouple. As a result, the size
difference is larger in the range where the coefficient of thermal
expansion is large (high-temperature side). To confirm this, in a
graph of the measurement results for a not-illustrated low-carbon
material (C: 0.002 wt %), both curves matched when displayed with
an offset of +4.degree. C. on the heating curve and -4.degree. C.
on the cooling curve. In the measurement of this diagram, a
stabilizing treatment at 98.degree. C. for 48 hours was
subsequently carried out, then the material was left at room
temperature for 816 hours until the measurement and the third cycle
measurement was carried out.
[0159] Similar to the measurement results of the first cycle, in
the third cycle a fairly large amount due to the seasoning effect
was again manifested. Compared with the 25 ppm in the first cycle,
this was 30 ppm in the third cycle. This is thought to be due to
the long standing time at room temperature.
[0160] The stabilizing temperature must be higher than room
temperature and be at least less than the Curie temperature.
However, since the actual Curie temperature should basically be at
the high-temperature stable position, a temperature of 150.degree.
C. or less, which is where the heating curve and the cooling curve
intersect, is realistic. In the present measurement, since it took
about 2 hours to heat from room temperature to 150.degree. C., it
can be said that there is an artificial seasoning effect even with
a treatment of 150.degree. C. for 2 hours or less.
[0161] Further, while it can be understood from the heating curve,
the temperature at which the heating curve and the cooling curve
(straight portion of 100.degree. C. or less) are no longer parallel
is 80.degree. C. At temperatures higher than that, these curves are
not parallel. From this result, it can be estimated that the carbon
atoms present in room temperature stable positions at 80.degree. C.
or more start to diffuse to high-temperature stable positions.
Therefore, to make most of the carbon atoms be in room temperature
stable positions, it is more desirable for the temperature to be
above room temperature to 80.degree. C. or less.
[0162] The fourth cycle followed the same trajectory as the second
cycle. Since measurement is carried out immediately after the third
cycle, and the stabilizing treatment is not carried out, this is a
projected result.
[0163] From the above, a member which has finished the second to
fourth cycles will have undergone the same treatment as gradually
cooling from a temperature of 205 to 210.degree. C. However, the
thus-treated alloy should not be fitted into a product. This is
because, for a high-carbon material, temporal deformation which is
larger only for 25 to 30 ppm than a member which has undergone a
stabilizing treatment will occur in the future.
[0164] Further, the fact that the high-carbon materials are subject
to temporal deformation of expanding by 18 ppm calculated on an
annual basis was described above. However, this measurement of
temporal deformation was carried out on a material after it had
already undergone a stabilizing treatment. Thus, this does not
contradict the fact that the same amount due to seasoning effect is
produced even for a material left at room temperature for just 1359
hours (about 2 months).
[0165] While the dimensional change of a low-carbon material test
piece was measured in the same manner as for the high-carbon
materials, regardless of whether the test piece was left at room
temperature before measurement or not, the sample length returned
to the initial size even after the heating and cooling.
Specifically, the contraction which was seen for the high-carbon
material could not be confirmed. Therefore, as mentioned above, the
reason why the slope becomes smaller at 80.degree. C. in the
heating curve of the high-carbon material and the size is
contracted at the point when cooling was carried out to room
temperature can be assumed to be due to the diffusion of carbon
atoms.
[0166] While the present invention has been described with
reference to exemplary embodiments, it is to be understood that the
invention is not limited to the disclosed exemplary embodiments.
The scope of the following claims is to be accorded the broadest
interpretation so as to encompass all such modifications and
equivalent structures and functions.
[0167] This application claims the benefit of Japanese Patent
Application Nos. 2008-117352, filed Apr. 28, 2008 and 2009-097226,
filed Apr. 13, 2009, which are hereby incorporated by reference in
their entirety.
* * * * *