The present invention relates to an advanced austenitic
stainless steel, more particularly to a structural austenitic stainless steel suited
for use in a corrosive environment or a high stress loaded environment, and to a
method for manufacturing the same.
Among the industrial steel materials, austenitic stainless
steel has widely been used as a structural material because of its excellent corrosion
resistance and workability. This steel, however, is low in strength in comparison
with other types of structural steel. Also, although the austenitic stainless steel
has high corrosion resistance, it is rather inferior to other types of steel in
use in a specific corrosive environment where pitting or stress corrosion cracking
is likely to occur.
With the progress of working efficiency and weight reduction
of the products in recent years, request has been rising in the industries for high
strength structural materials. In order to comply with such request, attempts have
been made to achieve higher strength of structural materials by use of additional
alloy elements such as rare metals, but use of such elements is not preferable in
view of recycling of the materials. One method for improving strength and corrosion
resistance of an alloy without changing the alloy composition is to utilize finer
The primary cause of deformation of metal material is slip
deformation caused by the so-called dislocation, which is the transfer of lattice
defect existing in the crystal. A high resistance is produced due to the interaction
between the grain boundaries and dislocation when dislocation passes across the
crystal grain boundaries.
To decrease grain size is to raise the density of the crystal
grain boundaries, and the phenomenon of deformation resistance increased by decreasing
grain size is well known as the Hall-Petch relationship, i.e., yield stress increases
in proportion to the -1/2 power of the crystal grain size.
For the alloys containing an element capable of forming
a protective film, such as Cr, the finer the crystal grains are, the more promoted
is the diffusion of the grain boundaries, whereby formation of the protective film
is made easier. The impurity element segregated at the grain boundaries is considered
as one of the causes of intergranular corrosion, but introduction of the grain boundaries
into bulk as a result of grain refinement may dilute the concentration of the impurity
element down to improve corrosion resistance. Conventional means for fining the
crystal grains of steel materials include thermomechanical treatment comprising
combination of such operation as rolling or upsetting with ensuing heat treatment.
Researchers are pursuing studies for comminuting the crystal
grains of austenitic stainless steel to the submicron size by thermomechanical treatment
making use of deformation-induced martensite transformation and inverse transformation
caused at a high temperature. Such studies are reported in, for instance,
Tetsu-to-Hagane, The Iron and Steel Institute of Japan, Vol. 80, pp. N529-N535,
Bulletin of Japan Institute of Metals, Vol. 27, No. 5, pp. 400-402, 1988
However, as is generally conceived, in a process where
a solid soluted material is rolled at a high draft, the crystal grain size is strongly
affected by workability, namely by the degree of working in the direction of rolling
and in the thickness direction, and tends to have a non-uniform distribution, so
that this method is unsuited for obtaining thick-walled components. Further, it
is not easy to obtain a high degree of working by cold rolling while avoiding the
formation of cracks.
Mechanical milling (mechanical alloying or mechanical grinding),
which performs forced working of metal powder by a ball mill or such, is capable
of forming powder having a crystal structure of nanometer grain size, since the
strain energy accumulated by working is much larger than the conventional methods
such as rolling. For consolidating the powder which has undergone mechanical milling
(hereinafter referred to as mechanically milled powder), the powder needs to be
sintered at a high temperature under a high pressure. Usually, strain energy is
released in the course of high-temperature heating to cause coarsening of the crystal
grains, so that it is difficult to carry out the consolidation process of the powder
while maintaining the nano-scale crystalline state.
Studies are underway for obtaining a bulk material of austenite
stainless steel by consolidating its mechanically milled powder and grinding the
crystal grains to the submicron size. They are reported, for instance, in (1)
, and (3)
Tetsu-to-Hagane, The Iron and Steel Institute of Japan, Vol. 84, pp. 357-362,
In the materials disclosed in (2) and (3) above, the sigma
phase is dispersed to control the growth of austenite crystal grains. However, the
M23C6 type carbide or sigma phase, which emerges principally
in austenitic stainless steel, is mainly composed of Cr, so that it acts to lower
Cr concentration in the surrounding and to encourage corrosion. It is possible to
reduce the influence of such carbide or sigma phase by reducing the grain size,
but such material can not be deemed suited as dispersed grains to be reduced in
The material described in the above literature (1) suggests
that precipitation of a carbide or oxide mainly composed of Ti, Zr or Nb is likely
to take place, but this literature fails to mention the optimal composition or process
conditions for controlling the grain growth.
Abstract of the Proceedings of 1998 General Meeting of The Iron and Steel
Institute of Japan, Vol. 11, page 563
, it is reported that the fine crystal structure of ferrite steel with
a nanometer grain size can be stabilized over a high temperature level of 1,000°C
or higher by adding and dispersing finely divided Y2O3 with
a grain size of several tens nanometers in the ferrite steel. However, when recycling
of the material is considered, addition of a specific alloy element such as yttrium
in the steel material is undesirable as it may complex the refining process, leading
to a rise of production cost.
As viewed above, with the techniques disclosed hitherto,
the manufacture of bulk material having a nona-scale ultra-fine crystal structure
is possible only under the conditions in which the dimensions and shape of the product
are restricted. Also, no disclosure has been made on the optimal compositions or
process conditions for achieving high strength and high corrosion resistance.
Accordingly, an object of the present invention is to provide
austenitic stainless steel of ultra-fine crystal structure having high strength
and high corrosion resistance in comparison with the conventional steel materials,
and a method for manufacturing such austenitic stainless steel.
The aspects of the present invention for attaining the
above object are as follows.
A corrosion resistant, high strength austenitic stainless
steel consisting of 1.0% or less of Si, 2.0% or less of Mn, 0.5% or less of O, 7
to 30% of Ni, 14 to 26% of Cr, 0.3% or less of combination of C and N, at least
one element selected from the group consisting of 1.0% or less of Ti, 2.0% or less
of Zr and 2.0% or less of Nb optionally 3.0% or less of Mo, and the balance consisting
of Fe and unavoidable impurities, the percentage being given in weight basis; said
steel containing carbonitride with a grain size of less than 100 nm dispersed therein;
said steel having an average crystal grain size of 1 µm or less; and
said steel containing 90% by volume or more of austenite phase.
The method for manufacturing a corrosion resistant, high
strength austenitic stainless steel, which comprises the steps of:
wherein said consolidating is carried out at 700 to 900°C. after the mechanically
processed powder has been retained at a temperature within a range of from 400 to
650 °C for a period of 0.5 to 6 hours, or alternatively, after the mechanically
processed powder has been suffered from a rise of temperature from 400 to 650 °C
for a period of 0.5 to 6 hours.
- providing a mechanically milled powder with an average crystal grain size of
200 nm or less consisting of 1.0% or less of Si, 2.0% or less of Mn, 0.5% or less
of O, 7 to 30% of Ni, 14 to 26% of Cr, 0.3% or less of combination of C and N, at
least one element selected from the group consisting of 1.0% or less of Ti, 2.0%
, or less of Zr and 2.0% or less of Nb, optionally 3.0% or less of Mo, and the balance
consisting of Fe and unavoidable impurities, the percentage being given in weight
- subjecting said mechanically processed powder to a process selected from the
group consisting of:
- (a) consolidating the mechanically milled powder at 700 to 900° C, and
- (b) consolidating the mechanically milled powder at 700 to 900°C to obtain
a consolidated material and thermomechanically treating the consolidated material
When two or more of Ti, Zr and Nb are used, the total amount
of the two or more of Ti, Zr and Nb is preferably 2.0% or less.
In the method for manufacturing a corrosion resistant,
high strength austenitic stainless steel the value f determined by the following
equation (1) falls within the range of from 0.4 to 2.0:
wherein (C), (N), (Ti), (Zr) and (Nb) are the amounts (weights) of the C, N, Ti,
Zr and Nb, respectively, in the mechanically milled powder.
In the method for manufacturing a corrosion resistant,
high strength austenitic stainless steel the value f determined by the following
equation (1) falls within the range of from 0.4 to 2.0:
wherein (C), (N), (Ti), (Zr) and (Nb) are the amounts (weights) of the C, N, Ti,
Zr and Nb, respectively, in the mechanically milled powder.
As said mechanically milled powder, it is possible to use
products obtained by subjecting a pre-alloy powder or a powder that meets the composition
defined above as a whole to mechanical grinding or alloying treatment with an attrition
mill or ball mill at 100°C or lower for 30 hours or more so that the products
have an average crystal grain size of 200 nm or less.
Preferably, said mechanical grinding or alloying treatment
with an attrition mill or ball mill is conducted using steel balls made of an Fe
alloy containing 0.3% or less of combination of C and N and having a heat conductivity
at 100°C of 16.7 W/m·K or higher.
Also preferably, said consolidation process are carried
out at 700 to 900°C after the mechanically milled powder has been retained
at a temperature within a range of from 400 to 650°C for a period of 0.5 to
6 hours, or alternatively, after the mechanically processed powder has been suffered
from a rise of temperature from 400 to 650°C for a period of 0.5 to 6 hours.
Preferably, the step of consolidating and the step of successive
thermomechanical treatment of the consolidated material include the step of consolidating
the mechanically milled powder by hot compression, hot rolling, hot isostatic pressing
or hot extrusion at 700 to 900°C, or the step of subjecting the consolidated
material to a heat treatment or hot forging at 700 to 900°C, and the additional
step of imparting a desired shape to the consolidated material during any of these
Said corrosion resistant, high strength austenite stainless
steel can be worked into a desired shape by press molding at 700 to 900°C.
Other objects, features and advantages of the invention
will become apparent from the following description of the embodiments of the invention
taken in conjunction with the accompanying drawings.
BRIEF DESCRIPTION OF THE DRAWINGS
- FIG. 1 is a schematic drawing illustrating the structure of a planetary ball
- FIG. 2 is a graph showing the influence of the added carbide-forming elements
on the strength of a hot isostatic pressing (HIP) worked material.
- FIG. 3 is microphotographs showing the metal structures of the Zr-added material
and the Zr-free material.
- FIG. 4 is a graph showing the relationship between the strength and the crystal
- FIG. 5 is a microphotograph and its schematic illustration showing the structure
of the material according to the present invention and the situation of distribution
of the carbonitride.
- FIG. 6 is a microphotograph and its schematic illustration showing the structure
of the non-added material and the situation of distribution of the carbonitride.
- FIG. 7 is a graph showing the influence of the temperature raising process on
the strength of the HIP worked material.
- FIG. 8 is a graph showing the balance of strength and toughness of the material
according to the present invention.
- FIG. 9 is anode polarization curves in a sulfuric acid solution.
- FIG. 10 is a perspective view showing the method of stress corrosion crack test.
- FIG. 11 is a schematic illustration of an attrition mill.
- FIG. 12 is a photograph showing the appearance of large-sized consolidated articles
made by hot isostatic pressing.
- FIG. 13 is a photograph showing the appearance of large-sized consolidated articles
made by hot forging a HIP worked material.
- FIG. 14 is a photograph showing the appearance of the consolidated articles
made by hot direct powder extrusion.
- FIG. 15 is a graph showing the relationship between consolidation temperature
vs. density of the consolidated articles.
- FIG. 16 is a photograph showing the appearance of a part shaped by hot press
The meaning of the numerals used in these drawings is as
DETAILED DESCRIPTION OF THE INVENTION
- 1: ball mill, 2: cover, 3: stock powder, 4: steel balls, 5: rotary table, 6:
mill holder, 7: fixture, 8: cooling fins, 9: crystal grains, 10: grain boundary,
11: inside of crystal grain, 12: Zr (C, N) and M23 (C, N)6
carbonitrides, 13: M23 (C, N)6 carbonitride, 14: plate test
piece, 15: grass fiber wool, 16: holder, 17: volt hole, 18: crushing tank, 19: cooling
water inlet, 20: cooling water outlet,
- 21: gas seal, 22: stock powder, 23: crushing steel balls, 24: agitator arm,
25: arm shaft.
The corrosion resistant, high strength austenitic stainless
steel according to the present invention is of a structure in which a carbonitride
with a grain size of several to 100 nm is dispersed, and has an average crystal
grain size is 1 µm or less and 90% by volume or more of austenite phase.
The method for manufacturing the steel according to the
present invention is characterized by including the steps of providing a mechanically
milled powder containing austenitic stainless steel main components, predetermined
amounts of carbonitride forming elements Ti, Zr and Nb, and C and N; vacuum sealing
the mechanically processed powder in a metallic container; and consolidating the
sealed powder at 700 to 900°C.
In the above method, the mechanically milled powder is
preferably the one prepared by subjecting a pre-alloy powder or a powder meeting
the above-defined composition as a whole to mechanical grinding or alloying treatment
by an attrition mill or ball mill at 100°C or lower for 30 hours or more so
that the powder will have an average crystal grain size of 200 nm or less, preferably
100 nm or less. Here, evaluation of crystal grain size is preferably made by microscopical
In the course of mechanical grinding or alloying treatment,
the elements such as Fe and Cr, in addition to C, N, O and H, may get into the powder
from the atmosphere, container, steel balls, stirring rod, etc., so that the chemical
composition of the austenitic stainless steel according to the present invention
is specified in relation to the mechanically milled powder and consolidated product.
In the present invention, especially C and N are the elements
which need to be controlled. The material of the steel balls, a main source of contamination,
is preferably made of Fe-base alloy that is equivalent of the mechanically milled
powder, which contains C and N in a combined amount of 0.3% or less, preferably
0.1 to 0.3%.
Also, for preventing accumulation of generated heat that
would cause overheating when conducting mechanical grinding or alloying treatment
on a large quantity of powder, it is preferable to use steel balls made of Fe alloy
having a heat conductivity at 100°C of 16.7 W/m·K or higher.
For obtaining the ultra-fine crystal structure of the austenitic
stainless steel according to the present invention, it is important to control the
grain growth during consolidating the mechanically milled powder having an average
crystal grain size of 200 nm or less, preferably 100 nm or less.
For controlling the grain growth by precipitating the carbonitride
and effectively pinning the intergranular transfer, it is necessary to carry out
the consolidating step after retaining the mechanically processed powder within
the temperature range of from 400 to 650°C for 0.5 to 6 hours, or after raising
the temperature to the range of 400 to 650°C over a period of 0.5 to 6 hours.
In case, for instance, a metastable austenitic stainless
steel composition corresponding to SUS304 is used as a base material, the composition
is transformed into a deformation-induced martensite structure when subjected to
mechanical grinding or alloying treatment, and for effecting inverse transformation,
it needs to heat the powder to 700°C or higher.
Generally, a higher temperature provides a higher sintering
efficiency, but since grain growth is accelerated at high temperatures, the temperature
should be kept 900°C or lower for obtaining a fine grain structure with a grain
size of 1 µm or less. Thus, the consolidating step is preferably carried out
at 700 to 900°C.
In the present invention, stainless steel having an austenite
structure is adopted as a base material. It is essential for the material according
to the present invention to be worked so as to have a mean grain size of 200 nm
or less, preferably 100 nm or less1, by mechanical grinding or alloying; but it
is not essential for the material to have a deformation-induced transformation structure
such as deformation-induced martensite structure.
However, mechanically milled powder having a deformation-induced
transformation structure is advantageous for forming finer grains, because, in the
heating process, the powder temporarily becomes finer than the pre-transformation
structure as a result of inverse transformation, whereby the effect of retarding
grain growth can be expected. Therefore, it is preferable to employ a chemical composition
that undergoes deformation-induced transformation when subjected to mechanical grinding
or alloying treatment.
By consolidating the mechanically milled powder by hot
pressing, hot rolling, hot isostatic pressing or hot extrusion at 700 to 900°C,
it is possible to compact the powder to substantially the same density as ingot
steels. Further, for improving ductility and toughness of the consolidated product,
it is desirable to carry out a heat treatment or hot forging at 700 to 900°C
on the consolidated product successively to the preceding step.
It is desirable to give the product a shape such as plate,
bar or a complicate shape in the course of the above steps for cost reduction via
process simplification. Also, for effecting shaping with a relatively low stress,
it is desirable to perform working at 700 to 900°C at which the super-plastic
mechanism of the fine structure is activated.
It is also desirable to give the product a desired shape
by press molding the steel material according to the present invention at 700 to
900°C for elongation of mold life or improvement of productivity.
The desirable form of structure for elevating strength
and corrosion resistance is the one in which the alloy has as its matrix an austenite
phase mainly composed of Fe-Cr-Ni or Fe-Cr-Ni-Mo, and the mean grain size of the
matrix-forming crystals falls within the range of from 30 to 1,000 nm. Since the
presence of other phases is detrimental to corrosion resistance, it is desirable
that the austenite phase accounts for at least 90% by volume of the structure.
Presence of a large number of grain boundaries increases
deformation resistance to enhance strength of the material. Also, impurities such
as P and S are segregated at the grain boundaries to make them the sites of corrosion,
but when they are introduced to a high density, the impurities are dispersed to
make the material highly resistant to local corrosion. Further, the grain boundaries
promote diffusion of the protective film forming elements such as Cr to the surface,
thus suppressing corrosion by rapid formation of protective film.
In the case of austenitic stainless steel, carbonitrides
such as M23 (C, N)6 and M (C, N) or sigma phase are formed
as fine precipitates that control grain growth in the consolidation process. The
carbonitride M (C, N) (M being Ti, Zr or Nb) can be precipitated at a relatively
low temperature in the structure where the defects have been introduced in large
quantities like the mechanically milled powder. The carbonitride is precipitated
finely, and its coarsening speed is low.
The M23 (C, N)6 carbonitride and
sigma phase are coarser than M (C, N) and composed mainly of Cr, so that they lower
the Cr concentration in the surrounding to give adverse effect to corrosion resistance.
Therefore, the M (C, N) carbonitride is suited as a grain growth inhibitory precipitate,
and its grain size is preferably in the range of several to 100 nm.
For preventing deterioration of corrosion resistance, it
is imperative to inhibit the formation of the M23 (C, N)6
carbonitride and sigma phase as much as possible. For this purpose, it is important
that Ti, Zr and Nb, which encourage the formation of sigma phase, be precipitated
as M (C, N) carbonitride before formation of sigma phase, and that C and N, which
forms the M23 (C, N)6 carbonitride, be precipitated as M (C,
N) carbonitride before precipitation of M23 (C, N)6 carbonitride.
Since Ti, Zr and Nb have high affinity for C and N, M (C,
N) carbonitride is more stable than M23 (C, N)6 carbonitride,
and in a structure where a large quantity of defects have been introduced, such
as the mechanically milled powder, it is possible to let M (C, N) carbonitride be
precipitated preferentially even at a relatively low temperature.
However, after the precipitation of M (C, N) carbonitride,
if the amount of C and N dissolved in the matrix is still high, a large quantity
of M23 (C, N)6 carbonitride will be formed, while if the amount
Ti, Zr and Nb dissolved in the matrix is still high, the sigma phase will be formed.
Therefore, the compositional ratios of Ti, Zr and Nb or C and N need to fall within
the range defined by the above-shown equation (1). It is also necessary to minimize
the amounts of Si and Mn, which promote the formation of sigma phase, and to fix
Si as an oxide.
Cr needs to be contained in an amount of 14% or more for
improving corrosion resistance. A high Cr content, however, destabilizes the austenite
phase and also encourages the formation of sigma phase to embrittle the material,
so that the Cr content is set at 26% at maximum, preferably 14 to 26%.
Mo is an additive element for raising corrosion resistance
and strengthening the solid solution. It is to be noted, however, that addition
of this element in excess of 3% extraordinarily promotes the formation of sigma
phase to embrittle the material, so that its content is preferably limited to 3%
or less for obtaining desired corrosion resistance and strength.
Ni has an action to stabilize the austenite phase to improve
corrosion resistance. A metastable composition which can cause deformation-induced
martensite transformation is advantageous for fining the crystal grains, but a low
Ni content is preferable, with its lower limit being 7%.
On the other hand, this element is usually contained in
an amount of 9% or more for raising corrosion resistance of the material. A high
Ni content is preferable for improving corrosion resistance, but when this element
is used in the same corrosive environment with other parts, it may trigger an electrochemical
reaction at the contact portion to promote corrosion of other parts, so that the
upper limit of its content is preferably set at 30%.
Ti, Zr and Nb, when added to a steel material, are usually
precipitated as M (C, N) carbonitrides to strengthen the material. They also have
a function to control crystal grain growth and serve for fining the crystal grains
of other M23 (C, N)6 carbides. In this powder-based alloy,
these elements act as a getter of oxygen impurities to purify the matrix.
On the other hand, excess addition of these elements to
the alloys results in embrittlement of the material. The preferred content of Ti,
when added, is 1.0% or less, and the preferred content of Zr and Nb is 2.0% or less.
In case two or more of Ti, Zr and Nb are added at the same
time, their total content is preferably restricted to 2% or less for controlling
excess precipitation of carbides. When the total content exceeds 2%, precipitation
of carbides increases to cause embrittlement of the material.
C and N are preferably contained in an amount of at least
0.02% for strengthening the solid solution and encouraging precpiptation of carbonitrides.
However, excess addition of these elements can give rise to excessive precipitation
of chromium carbonitride to cause lowering of corrosion resistance due to the decrease
of the amount of chromium dissolved in the matrix. Therefore, the upper limit of
the content of these elements is preferably set at 0.3%.
Oxygen (O) is already contained as an impurity in the preparation
of the powder. It also gets mixed in the material in the step of mechanical milling.
This element produces oxides to strengthen the material, but the oxide film on the
powder surface impedes sintering, and excessive formation of oxides leads to embrittlement
of the material, so that the content of O is set at 0.5% by weight at maximum, preferably
0.1% by weight or less. As this element has the role of fixing sigma phase-forming
Si as an oxide, its content is preferably decided depending on the amount of Si
Si and Mn are added as a deoxidizer in the preparation
of the powder, Mn also serving as a desulfurizer. Si is eluted out from the melting
crucible in the powder preparation and mixed in the material. Since Si and Mn promote
the formation of sigma phase, their content is preferably lessened as much as it
can be. Si is contained in an amount of 1.0% or less, preferably 0.6% or less, and
Mn in an amount of 2.0% or less, preferably 0.2% or less, according to the JIS standards
of austenitic stainless steel.
P and S, which have an action to reduce corrosion resistance,
are contained in the course of preparation of the powder. Preferably P is contained
in an amount of 0.045% or less and S in an amount of 0.030% or less according to
the JIS standards of austenitic stainless steel.
First, a method for manufacturing a nano-size crystalline
steel material according to the present invention is described. In this example,
a planetary ball mill illustrated in FIG. 1 was used for mechanical grinding and
160 g of starting powder 3 and 9.5 mm diameter stainless
steel balls 4 were filled in a stainless steel-made 470 ml-capacity ball mill 1
containing argon gas, and the mill was closed airtight by a stainless steel cover
2. The packed ball mill 1 was secured to a mill holder 6 on a ball mill rotary table
5 by a fixture 7. Rotating force was transmitted to the rotary table 5 from an outside
driving system, producing a centrifugal force in four sets of ball mill 1 disposed
crosswise on the table, each ball mill rotating itself its own axis, causing the
steel balls 4 to collide against each other or against the inner wall of the mill
1. Consequently, the starting powder 3 was forcibly worked to produce a mechanically
processed powder having a large quantity of defects and fine crystal grains with
an average size of around 50 to 100 nm. The speed of the rotating table 5 was set
at 200 rpm.
Starting powder was a pre-alloy powder or a mixed powder
conforming to the specified composition as a whole. It was possible to obtain the
mechanically processed powder in each case by adjusting the milling conditions.
Ball mill 1 was cooled by the cooling fins 8 provided in the mill and by the blast
produced by the rotation. It was confirmed that the temperature was maintained at
50°C during milling upon measuring the temperature immediately after milling.
The main chemical components (wt%) of each type of austenitic
stainless steels with fined crystal grains are shown in Table 1.
The mechanically milled powder was contained in the vacuum
sealed, mild steel-made capsules and subjected to hot isostatic pressing in 196
MPa of argon gas at 800 to 900°C for more than one hour to obtain a consolidated
product compacted to substantially the same density as ingot steel of the same composition.
FIG. 2 shows difference in strength of the consolidated
products immediately after hot isostatic pressing, with the type of the carbide-forming
additive element Zr, Ti, Nb or Mo.
Clearly, the Zr-, Ti- and Nb-added materials (Nos. 6, 7
and 8, respectively, in Table 1) are improved in strength in comparison with the
0.12% C-added material (No. 5 in Table 1).
Table 2 shows the carbonitrides identified by X-ray diffractometry.
It is seen that M23 (C, N)6 carbonitride alone is observed
in the 0.12% C-added material (No.5 in Table 1) and in the V- and Mo-added materials
(Nos. 9 and 10, respectively, in Table 1), whereas M23 (C, N)6
carbonitride as well as M (C, N) carbonitride are observed in the Zr-, Ti- and Nb-added
materials (No. 6, 7 and 8, respectively, in Table 1). This shows that the presence
of M (C, N) carbonitride contributes to the improvement of strength.
As a representative example, the structure of the Zr-added
material (No. 6 in Table 1) is shown in FIG. 3. Clearly, the crystal grains 9 in
this material are finer than the non-added material (No. 1 in Table 1).
Carbonitrides identified by X-ray diffractometry
Ti(C,N), M23(C, N)6
The relationship between strength and crystal grain size
shown in FIG. 4 is substantially in accord with Hall-Petch relationship, which indicates
that grain refinement is the main factor of strengthening.
Closer observation of the structure revealed dispersion
of fine particles of Zr (C, N) and M23 (C, N)6 carbonitrides
12 at the crystal grain boundaries 10 and in the inside 11 of the crystal grains
as seen in FIG. 5. The grain size of the carbonitride in the crystal grains was
several to several tens nm, and that of the carbonitride at the crystal grain boundaries
was several tens to 100 nm.
On the other hand, the microscopic structure of the non-added
material (No. 1 in Table 1) shows that M23 (C, N)6 carbonitride
13 was dispersed at the crystal grain boundaries 10 and in the inside 11 of the
crystal grains as shown in FIG. 6, the grain size of the carbonitride in the crystal
grains being several tens to 100 nm and that of the carbonitride at the crystal
grain boundaries being 100 to 200 nm.
As described above, by the addition of Zr, Ti or Nb, the
fine grains of M (C ,N) and M23 (C, N)6 carbonitrides 10 such
as shown in FIG. 5 are precipitated and dispersed to pin intergranular transfer,
with the result that the grain growth during the consolidation process is controlled
and a finer structure can be obtained.
In the heating process during consolidating, the temperature
at which the precipitation of carbonitride begins is lower than the temperature
at which the grain growth is promoted. The temperatures from 400 to 650°C are
the temperature zone in which the grain growth is not promoted but the carbonitrides
are precipitated, so that by effecting sufficient precipitation of the carbonitrides
by keeping the temperature within the range, it is possible to control the grain
growth in the ensuing high temperature process to provide high strength.
Regarding, for instance, Zr-added material (No. 11 in Table
1) and Ti-added material (No. 15 in Table 1), when these materials are once held
in the temperature zone of 500 to 650°C in the heating step in the hot isostatic
pressing process, a peak of strength is seen as shown in FIG. 7. In the case of
these materials, the temperature was raised to and maintained at 400°C in the
vacuum sealing step before the hot isostatic pressing, and with the material which
was not held in the specified temperature range in the heating step in the hot isostatic
pressing process, data were plotted at 400°C in FIG. 7.
Table 3 shows the mechanical properties of the consolidated
products subjected to hot isostatic pressing process and then hot forging at 850°C.
Average crystal grain size (pm)
0.2% proof strength (MPa)
Tensile strength (MPa)
a: materials of the present
invention b: comparative materials
b*: comparative materials (ingots)
Ductility was remarkably improved by hot forging. A diagrammatic
illustration of comparison with the conventional materials regarding toughness and
tensile strength is shown in FIG. 8. Toughness was evaluated by absorbed energy
determined by the V-notched Charpy impact test pieces.
Conventional material (indicated by shadowed rhombi in
FIG. 8) was a solid soluted material of austenitic stainless steel with a composition
of 18% Cr and close to 8% Ni (compositions of Nos. 22 to 26 in Table 1); Conventional
material 2 (indicated by unshadowed rhomb in FIG. 8) was a cold worked material
of austenitic stainless steel (composition of No. 22 in Table 1); and Conventional
material 3 (indicated by shadowed triangle) was semi-austenitic precipitation hardened
stainless steel (composition of No. 27 in Table 1).
The conventional materials would prove low in toughness
if they were provided with high strength, but the steel material according to the
present invention is high in both strength and toughness. Some conventional materials
are within the same range of composition as the material according to the present
invention, but they are low in strength because of their coarse structure due to
the difference in the manufacturing process. Although the conventional material
of No. 1 (comparative material) has an improved strength by cold working, it is
low in toughness.
FIG. 9 shows the result of evaluation of corrosion resistance
of the material according to the present invention by determination of anode polarization
curve in a sulfuric acid solution (1N, 30°C).
The material according to the present invention (for example,
No. 7 in Table 1) is low in critical passivation current density and passive state
maintaining current density in comparison with the conventional materials. This
indicates that the material according to the present invention exhibits higher corrosion
resistance than the conventional materials at a potential in or lower than the passive
Stress corrosion crack resistance of the material according
to the present invention was evaluated by a CBB test. A perspective view of the
testing apparatus is shown in FIG. 10.
A plate-like test piece 14 was held between rounded holders
16 together with a spacing piece of glass fiber wool 15 and bolts were passed through
bolt holes 17 and fastened. This assembly was immersed in an autoclave filled with
high-temperature (288°C) and high-pressure (85 kg/cm2) pure water
(with 8 ppm of oxygen dissolved therein). The immersion time was 500 hours. After
the test, each test piece was checked for cracks by its sectional observation under
a light microscope. As a result, no crack was observed in all of the test pieces
of the materials according to the present invention.
In this example, an attrition mill shown in FIG. 11 was
used for mechanical milling. This attrition mill comprised a stainless crushing
tank 18 with a capacity of 25 liters, a cooling water inlet 19 of the tank 18, a
cooling water outlet 20, a gas seal 21 for sealing a replacing gas such as argon
or nitrogen gas, 10 mm-diameter crushing steel balls 23 placed in the tank, and
an agitator arm 24. 5 kg of starting powder 22 was contained in the tank.
A pre-alloy powder corresponding to No. 11 in Table 1 was
used as starting powder 22. Rotating force was transmitted to arm shaft 25 from
an outside driving system to rotate agitator arm 24. Steel balls 23 were agitated
by said rotating agitator arm 24 so that these steel balls 23 collide against each
other or against the inner wall of tank 18, whereby mixed powder 22 was forcibly
processed to produce an alloy powder composed of fine crystal grains.
The speed of arm shaft 25 was set at 150 rpm, and the treating
time was 50 hours.
Regarding the powder subjected to mechanical milling by
the planetary ball mill in Example 1 or the attrition mill in the instant example,
the increments of carbon concentration and the type of steel balls used are shown
in Table 4.
Type of steel balls
Properties of steel
Increment of carbon
Carbon concentration (wt%)
Heat conductivity (W/(m·K))
Planetary ball mill
Amount treated: 160 g
Amount treated: 5 kg
Untreatable because of overheating
When using the steel balls B with a high carbon concentration
of 1.05% by weight, the carbon concentration in the processed powder rises up greatly.
For the material according to the present invention, it is necessary to strictly
control the carbon concentration, and any excess increase of carbon concentration
during mechanical grinding is undesirable.
In case of using the steel balls A having substantially
the same composition as starting powder, although no problem arises when a small
quantity of powder is treated by a planetary ball mill, the concern is that the
inside of the mill may be overheated, making it unable to continue the treating
operation, when a large quantity of powder is treated by an attrition mill.
According to the present invention, it was possible to
solve the problems of increase of carbon concentration and overheating in the inside
of the mill by using the steel balls C with high heat conductivity and a carbon
concentration of the same level as the material according to the present invention.
20 kg of mechanically milled powder, prepared by conducting
mechanical grinding by an attrition mill, was vacuum sealed in the mild steel-made
capsules and then subjected to consolidation process by hot isostatic pressing.
In the heating step of the hot isostatic pressing process, the powder was once held
at around 600°C for one hour and then at 850°C in 196 MPa argon gas for
3 hours. As a result, there could be obtained 20 kg of a consolidated product such
as shown in FIG. 12.
A tensile test was conducted on the test pieces cut out
from the various parts of this consolidated product to evaluate its homogeneity.
It was confirmed that the respective parts of the consolidated product were almost
equal in durability and had tensile ductility, that is, the product has been sintered
Next, 20 kg of a consolidated product made by hot isostatic
pressing was heated to 850°C and then forged to an upsetting ratio of about
3. As a result, it was possible to effect desired deformation with no problem as
shown in FIG. 13.
The result of the tensile test conducted on the cut out
test pieces confirmed improvement of tensile ductility by hot forging, as same as
shown in Table 3 in Example 1. It was also confirmed that the product shows high
strength without decreasing Charpy absorption energy as in the case of FIG. 6 in
2.8 kg of a mechanically milled powder produced according
to the process of Example 2 was vacuum sealed in the mild steel capsules and then
consolidated by hot direct powder extrusion. In the extrusion process, the powder
was once held at close to 600°C for 2 hours and then at 750°C, 800°C
and 850°C for one hour each, and thereafter extruded from a die set at an extrusion
ratio of 5.7.
As a result, it was possible to obtain the bars conforming
to the appearance requirements as shown in FIG. 14. A tensile test was conducted
on the cut out test pieces, confirming that it was possible to obtain a tensile
strength of 1,000 MPa or higher and a tensile ductility of 30% or higher.
The relation of the consolidation temperature to the density
of the consolidated product is shown in FIG. 15. In the case of the consolidated
product obtained by hot isostatic pressing of 196 MPa, it was necessary to raise
the consolidation temperature to 800°C or above for compacting the product
to substantially the same density as the ingot steel, but in the case of the consolidated
product made by hot extrusion, it was possible to obtain substantially the same
density as the ingot steel even at a consolidating temperature of 750°C.
The consolidated product made in Example 3 was cut into
a disc, and the disc was heated to 750°C, then held in a heated mold and compressed
by a press. It was confirmed that a part of such complicate configuration as shown
in FIG. 16 could be produced with a lower compressive force than required for ingot
According to the present invention, a bulk material of
austenitic stainless steel can be obtained by consolidating a powder having its
structure ultra-fined by mechanical milling while controlling the crystal grain
growth. Such bulk material has higher strength and higher corrosion resistance with
less reduction of toughness than that obtainable with the conventional methods,
and in such bulk material, the nano scale fine crystal grain structure is uniformly
distributed and its properties are uniformalized throughout the whole material.
Further, according to the present invention, it is possible
to provide a molded article with a complicate configuration by hot forging with
lower stress than required for the conventional materials.
Since a nona scale fine crystal grain structure can be
obtained with a composition comprising the alloy elements contained in the ordinary
steel materials, it is possible to provide austenitic stainless steel with excellent
It should be further understood by those skilled in the
art that the foregoing description has been made on embodiments of the invention
and that various changes and modifications may be made in the invention without
departing from the spirit of the invention and the scope of the appended claims.