U.S. patent application number 10/987878 was filed with the patent office on 2006-02-02 for ultratough high-strength weldable plate steel.
This patent application is currently assigned to Northwestern University. Invention is credited to Gregory B. Olson, Arup Saha.
Application Number | 20060021682 10/987878 |
Document ID | / |
Family ID | 35197564 |
Filed Date | 2006-02-02 |
United States Patent
Application |
20060021682 |
Kind Code |
A1 |
Saha; Arup ; et al. |
February 2, 2006 |
Ultratough high-strength weldable plate steel
Abstract
A transformation toughened, high-strength steel alloy useful in
plate steel applications achieves extreme fracture toughness
(C.sub.v>80 ft-lbs corresponding to K.sub.Id.gtoreq.200
ksi.in.sup.1/2) at strength levels of 150-180 ksi yield strength,
is weldable and formable. The alloy is characterized by dispersed
austenite stabilization for transformation toughening to a
weldable, bainitic plate steel and is strengthened by precipitation
of M.sub.2C carbides in combination with copper and nickel. The
desired microstructure is a matrix containing a bainite-martensite
mix, BCC copper and M.sub.2C carbide particles for strengthening
with a fine dispersion of optimum stability austenite for
transformation toughening. The bainite-martensite mix is formed by
air-cooling from solution treatment temperature and subsequent
aging at secondary hardening temperatures to precipitate the
toughening and strengthening dispersions.
Inventors: |
Saha; Arup; (Hillsboro,
OR) ; Olson; Gregory B.; (Riverwoods, IL) |
Correspondence
Address: |
BANNER & WITCOFF, LTD.
TEN SOUTH WACKER DRIVE
SUITE 3000
CHICAGO
IL
60606
US
|
Assignee: |
Northwestern University
Evanston
IL
|
Family ID: |
35197564 |
Appl. No.: |
10/987878 |
Filed: |
November 12, 2004 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
60519388 |
Nov 12, 2003 |
|
|
|
Current U.S.
Class: |
148/547 ;
420/91 |
Current CPC
Class: |
C21D 2211/008 20130101;
C21D 2211/002 20130101; C21D 2211/004 20130101; C21D 6/004
20130101; C22C 38/46 20130101; C21D 6/02 20130101; C21D 1/20
20130101; C21D 1/78 20130101; C22C 38/42 20130101; C22C 38/44
20130101; C21D 2211/001 20130101; C21D 1/25 20130101; C21D 8/0263
20130101 |
Class at
Publication: |
148/547 ;
420/091 |
International
Class: |
C22C 38/42 20060101
C22C038/42; C22C 38/44 20060101 C22C038/44 |
Goverment Interests
REFERENCE TO GOVERNMENT RESEARCH CONTRACTS
[0002] This development was supported by the Office of Naval
Research (Grant No. N00014-01-1-0953.
Claims
1. A weldable, workable steel comprising, in combination: a
secondary hardened, stable, generally mixed bainite and martensite
phase steel having a BCC copper precipitation phase and M.sub.2C
carbide strengthening particles, said steel including in weight
percent about: 0.30 to 0.55 carbon (C); 3.5 to 5.0 copper (Cu); 6.0
to 7.5 nickel (Ni); 1.6 to 2.0 chromium (Cr); 0.2 to 0.6 molybdenum
(Mo); 0.05 to 0.20 vanadium (V); and the balance Fe.
2. The steel of claim 1 characterized by a yield strength exceeding
about 140 ksi.
3. A steel having enhanced toughness, weldability strength and
workability, said steel comprising, in combination in weight
percent about: 0.30 to 0.55 carbon (C); 3.50 to 5.0 copper (Cu);
6.0 to 7.5 nickel (Ni); 1.6 to 2.0 chromium (Cr); 0.2 to 0.6
molybdenum (Mo); 0.05 to 0.20 vanadium (V), and the balance iron
(Fe) characterized by a yield strength exceeding about 140 ksi, a
generally martensitic microstructure, BCC copper precipitate
hardening, M.sub.2C carbide particle strengthening where M is one
or more materials selected from the group consisting of Cr, Mo and
V, and Ni stabilized, precipitated austenite.
4. The steel of claim 3 wherein the steel is characterized by a a
mixed bainite and martensite microstructure.
5. The steel of claim 3 wherein the steel is characterized by
secondary hardening.
6. The steel of claim 3 characterized by a first tempering step in
the range of about 450.degree. C. for a period of about 30 to 90
minutes.
7. The steel of claim 5 characterized by a second tempering step in
the range of about 300.degree. C. to 450.degree. C. for a period of
about 1 to 10 hours.
8. The steel of claim 1 having a nominal composition in weight
percent of about: 0.05.+-.0.01 C; 3.65.+-.0.05 Cu; 6.5.+-.0.02 Ni;
1.84.+-.0.05 Cr; 0.60.+-.0.05 Mo; 0.10.+-.0.01 V; and the balance
Fe.
9. A method for manufacture of a hardened steel alloy comprising
the steps of: (a) forming a melt into a casting comprised in weight
percent of about 0.3 to 0.55 carbon (C), 3.5 to 5.0 copper (Cu),
6.0 to 7.5 nickel (Ni), 1.6 to 2.0 chromium (Cr), 0.2 to 0.6
molybdenum (Mo), 0.05 to 0.20 vanadium (V) and the balance iron;
(b) homogenizing the casting at a temperature in the range of about
1200.degree. C..+-.50.degree. for about 6 to 8 hours; (c) hot
working said alloy; (d) ambient cooling said alloy; (e) annealing
said alloy at a temperature in the range of about 480.degree.
C..+-.40.degree. C. for about 8 to 12 hours; (f) solution heating
said alloy at a temperature of about 900.degree. C..+-.50.degree.
C. for about 30 to 90 minutes; (g) cooling said alloy to form a
generally martensitic microstructure; (h) tempering said alloy at a
temperature of about 500.degree. C. to 575.degree. C. for about 5
to 90 minutes to achieve austenite precipitation; and (i) further
tempering said alloy at a temperature of about 400.degree. C. to
500.degree. C. for about 1 to 10 hours to achieve M.sub.2C carbide
particle formation.
10. A method for manufacture of a hardened steel alloy comprising
the steps of: (a) forming a melt into a casting comprised in weight
percent of about 0.3 to 0.55 carbon (C), 3.5 to 5.0 copper (Cu),
6.0 to 7.5 nickel (Ni), 1.6 to 2.0 chromium (Cr), 0.2 to 0.6
molybdenum (Mo), 0.05 to 0.20 vanadium (V) and the balance iron;
(b) heat treating and working said alloy to form a material
characterized by a substantially martentsitic microstructure; (c)
tempering said alloy at a temperature of about 500.degree. C. to
575.degree. C. for about 5 to 90 minutes to achieve dispersed
austenite precipitation; and (d) further tempering said alloy at a
temperature of about 400.degree. C. to 500.degree. C. for about 1
to 10 hours to achieve M.sub.2C carbide particle formation where M
is selected from the group consisting of Cr, V and Mo.
11. A method for manufacture of a hardened steel alloy comprising
the steps of: (a) forming a melt into a casting comprised in weight
percent of about 0.3 to 0.55 carbon (C), 3.5 to 5.0 copper (Cu),
6.0 to 7.5 nickel (Ni), 1.6 to 2.0 chromium (Cr), 0.2 to 0.6
molybdenum (Mo), 0.05 to 0.20 vanadium (V) and the balance iron;
(b) working and heat treating said alloy to form a generally
bainitic and martensitic microstructure; (c) tempering said alloy
at a temperature of about 500.degree. C. to 575.degree. C. for
about 5 to 90 minutes to achieve austenite dispersion
precipitation; and (d) further tempering said alloy at a
temperature of about 400.degree. C. to 500.degree. C. for about 1
to 10 hours to achieve M.sub.2C carbide particle formation wherein
M is selected from the group consisting of Cr, Mo and V, said alloy
including BCC Cu, Ni dispersed austenite and M.sub.2C carbides in a
mixed bainite and martensite microstructure.
Description
CROSS-REFERENCE TO RELATED APPLICATION
[0001] This is a utility application based upon and incorporating
by reference provisional application Ser. No. 60/519,388 filed Nov.
12, 2003 entitled Ultratough High-Strength Weldable Plate Steel for
which priority is claimed.
BACKGROUND OF THE INVENTION
[0003] In a principal aspect, the present invention relates to a
steel alloy and a process for making such an alloy which exhibits
new levels of strength and toughness while meeting processability
requirements. The ultratough, weldable secondary hardened plate
steel alloys for structural applications exhibits fracture
toughness (K.sub.Id.gtoreq.200 ksi.in.sup.1/2) at strength levels
of 150-180 ksi yield strength, is weldable and formable.
[0004] Throughout the history of materials development, there has
been an ever-increasing need for stronger, tougher, more fracture
resistant and easily weldable plate steels for structural
applications at minimal cost. Unfortunately, however, any increase
in strength is rarely achieved without concomitant decreases in
toughness and ductility, which limits the utility of most
ultrahigh-strength steels. The best combinations of strength and
toughness have usually been obtained from martensitic
microstructures as shown in FIG. 1.
[0005] High strength bainitic steels have not been as successful in
practice because of coarse cementite particles in bainite that are
detrimental to toughness. Nonetheless, a potential benefit
motivating research of air-hardened steels containing
bainite/martensite mixtures is the ease of processing, which may
lead to a product with good performance at a relatively lower cost.
The possibility of improving the strength and toughness
simultaneously using fine-grained bainitic ferrite plates and
enhancing the toughness by transformation toughening effects
presents a technological challange. Further improvements of
strength can possibly be achieved with co-precipitation of alloy
carbides and bcc copper for easily weldable, low-carbon steels
again presenting a technological challenge.
[0006] It is now known that the interaction of deformation-induced
martensitic transformation of dispersed austenite with
fracture-controlling processes such as microvoid induced shear
localization results in substantial improvements in fracture
toughness called Dispersed Phase Transformation Toughening (DPTT).
Transformation toughening is attributed to modification of the
constitutive behavior of the matrix through pressure-sensitive
strain hardening associated with the transformation volume change.
The transformation behavior and the toughening effects are
controlled by the stability of the austenite dispersion. For
transformation toughening at high strength levels, the required
stability of the austenite dispersion is quite high and can be
achieved only by size refinement and compositional enrichment of
the austenite particles. The size influences the characteristic
potency of nucleation sites in the particles while the composition
influences the chemical driving force and interfacial friction for
the martensitic transformation. The size refinement and the
compositional enrichment of the austenite can possibly be
controlled with heat treatments such as multi-step tempering.
[0007] With this general background, design objectives motivating
the invention are the achievement of extreme impact fracture
toughness (C.sub.v>85 ft-lbs corresponding to fracture
toughness, K.sub.Id>200 ksi.in.sup.1/2 and K.sub.Ic>250
ksi.in.sup.1/2) at high strength levels of 150-180 ksi yield
strength in weldable, formable plate steels with high resistance to
hydrogen stress corrosion cracking (K.sub.ISCC/K.sub.IC>0.5).
Design goals are marked by the star in the cross-plot of K.sub.Ic
fracture toughness and yield strength illustrated in FIG. 2. This
design aims to substantially expand the envelope marked as "steels"
to the top right corner of the plot. Optimization of such a system
and achievement of design goals can possibly be effectively
achieved by consideration of the methods of systems design. FIG. 3
describes in general a system approach to design steel with the
specified strength, toughness levels as well as optimum weldability
and hydrogen resistance.
[0008] As further background, recent studies have shown that
selection of fine Ti(C,N) as a grain refining dispersion
contributes to increasing the fracture resistance by delaying the
coalescence of microvoids among the primary voids. Studies have
also suggested that the resistance to primary void formation and
coalescence is proportional to inclusion spacing. Thus, it may be
desirable to reduce the volume fraction of primary inclusions or
coarsen inclusions for a given volume fraction. This can be
achieved by clean melt practices and tight composition control.
However, engineering design fracture toughness parameters like
K.sub.Ic and K.sub.Id are difficult and expensive to measure. Thus
for preliminary design analyses, small-scale inexpensive fracture
measurements like Charpy V-notch impact energy (C.sub.v) values may
be used to estimate K.sub.Ic and K.sub.Id. Studies of fracture
toughness dependence on loading rate measured over a temperature
range have shown that K.sub.Ic fracture toughness values under
static and intermediate loading are about 20% higher than the
K.sub.Id measured under impact loading. An approximate correlation
between K.sub.Ic and C.sub.v test results for conventionally
grain-refined steels is as follows: K.sub.IC.sup.2=AC.sub.v (1)
where A is a constant of proportionality. Fitting equation (1) to
results from high Ni steels is shown in FIG. 4.
[0009] According to these relationships, the C.sub.v impact
toughness objective of 85 ft-lbs corresponds to a K.sub.Ic fracture
toughness under static loading of 250 ksi.in.sup.1/2 and a dynamic
K.sub.Id of 200 ksi.in.sup.1/2.
[0010] A fine carbide dispersion may need to be obtained in order
to achieve the desired strength level. Coherent M.sub.2C carbides
have been used in secondary hardened steels that are currently in
use. Previous work to optimize the carbide particle size for
maximizing the strength 3 nm carbide precipitates corresponding to
the transition from particle shear to Orowan bypass may provide
maximum strength. Thermodynamics and kinetics of carbide
precipitation may need to be controlled to obtain such a fine
M.sub.2C carbide dispersion. The driving force for M.sub.2C
nucleation may also be maximized by proper control of the amount
and ratio of carbide formers in the alloy to refine the M.sub.2C
particle size. Sufficient M.sub.2C precipitation may need to be
achieved to dissolve cementite in order to attain the desired
toughness levels because coarse cementite particles are extremely
deleterious as microvoid nucleation sites. Tempering times should
also be minimized to prevent impurity segregation at grain
boundaries.
[0011] Even if low alloy carbon levels are maintained, steels
containing higher alloying content might help in achieving the
desired combination of mechanical properties, but may reduce the
weldability of the material by increasing hardenability. For any
structural material, the heat-affected zones (HAZ) adjacent to the
welded joints are considered to be the weakest links. Weldability
of steels is generally controlled by both the matrix and the
strengthening dispersion structures. As a rule of thumb, for
adequate weldability of the steel C content of the alloy should be
kept below 0.15 wt %. This in turn limits the C available for
M.sub.2C strengthening. For a bainitic matrix, modification of the
hardenability of the steel may provide bainite with a much lower
cooling rate. However, weldability can deteriorate as the
hardenability increases. Again, numerous interrelated technological
challenges are apparent in view of various known
considerations.
[0012] Ultra-high strength steels are prone to a decrease of
fracture toughness in aqueous environments due to hydrogen assisted
cracking. This reduction of toughness is caused by intergranular
brittle fracture associated with impurity segregation to grain
boundaries, which may reduce toughness of the steel by as much as
80% in a corrosive environment. The common impurities in steel are
P and S, both of which are embrittlers since they have lower free
energy on a surface than at a grain boundary. An effective way of
reducing them is by cleaner processing techniques or impurity
gettering. Impurity gettering can tie up P and S as stable
compounds formed during solidification. La and Zr have been found
to be effective impurity gettering elements. Another approach to
minimize impurity effects is by design of grain boundary chemistry.
Segregating elements like W and Re preferentially on the grain
boundaries that may enhance grain boundary cohesion could be
beneficial to the stress corrosion cracking resistance. Small
amounts of dissolved B may also help in grain boundary cohesion. In
view of the numerous foregoing factors and information, a need for
an improved high strength plate steel was addressed.
SUMMARY OF THE INVENTION
[0013] A transformation toughened ultratough high-strength steel
alloy useful in plate steel applications achieves extreme fracture
toughness (C.sub.v>80 ft-lbs corresponding to
K.sub.Id.gtoreq.200 ksi.in.sup.1/2) at strength levels of 150-180
ksi yield strength, is weldable and formable. The alloy employs
dispersed austenite stabilization for transformation toughening to
a weldable, bainitic plate steel and is strengthened by
precipitation of M.sub.2C carbides in combination with copper and
nickel. The desired microstructure is a matrix containing a
bainite-martensite mix, BCC copper and M.sub.2C carbides for
strengthening with a fine dispersion of optimum stability austenite
for transformation toughening. The bainite-martensite mix is formed
by air-cooling from solution treatment temperature and subsequent
aging at secondary hardening temperatures to precipitate the
toughening and strengthening dispersions.
[0014] More specifically, steel alloys nominally in weight percent
comprised of about 0.3 to 0.55 carbon (C), 3.5 to 5.0 copper, 6.0
to 7.5 nickel (Ni), 1.6 to 2.0 chronimym (Cr), 0.2 to 0.6
martensite (Mo), 90.05 to 0.20 vanadium (V) and the balance iron
(Fe) and insubstantial impurities is formed from a melt and heat
treated by various steps including tempering, for example, to form
an essentially banite/martensite phase alloy with dispersed
austenite, M.sub.2C carbide strengthening where M is Cr, V and/or
Mo, dispersed BCC copper for precipitation strengthening and nickel
to promote austenite stability and transformation toughening.
[0015] The solidified melt is preferably subjected to a two stage
tempering process with the first stage at a higher temperature in
the range of .degree. C. to 600.degree. for less than one hour
followed by a lower temperature stage for more that one hour of
about 400.degree. C. to 500.degree. C.
BRIEF DESRIPTION OF THE DRAWINGS
[0016] In this application reference is made to the drawing
comprised of the following figures:
[0017] FIG. 1 is a graph of K.sub.IC toughness vs. R.sub.C hardness
cross-plot for ultra-high strength martensitic steels.
[0018] FIG. 2 is a graph of K.sub.IC toughness vs. .sigma. yield
strength cross-plot for different classes of materials.
[0019] FIG. 3 is a systems design chart for blast resistant naval
hull steels.
[0020] FIG. 4 is a graph correlation between K.sub.Ic and C.sub.v
test results for high Ni steels.
[0021] FIG. 5 is a schematic of the design optimization
procedure.
[0022] FIG. 6 is a graph of power-law relationship relating
hardness of related steels to yield stress from experimental data
from Foley (circles), Kuehmann (triangles) and Spaulding (diamonds)
shown in comparison to straight-line relationship for ideal plastic
material.
[0023] FIG. 7 is a design Graville diagram for determining
susceptibility to HAZ cracking in plate steels.
[0024] FIG. 8 is a graph change in hardness as a function of alloy
carbon content for M.sub.2C carbide strengthening contribution. The
arrows represent hardness increment of 175 VHN is achieved at C
level of 0.05 wt % set for the alloy. Experimental results of other
secondary hardening steels are shown.
[0025] FIG. 9 is a graphical representation for contributions of
the individual mechanisms to achieve the strength goal equivalent
to 389 VHN.
[0026] FIG. 10 is a Cr--Mo Phase Diagram Section at 900.degree. C.
with alloy composition in atomic %:
Fe--0.234C--1.32Cu--6.21Ni--0.055V. This diagram shows the phase
fields of the FCC austenite and FCC+M.sub.6C revealing that the
M.sub.2C stoichiometric line is well within the solubility
limit.
[0027] FIG. 11 is a graph depicting driving Forces (in kJ/mole) for
M.sub.2C carbide nucleation contour plot varying at % (Mo) and at %
(Cr) with superimposed M.sub.2C stoichiometric (heavy) line at
500.degree. C. at alloy compositions at %
Fe--0.234C--1.32Cu--6.21Ni--0.055V.
[0028] FIG. 12 is a graph depicting driving Force (in kJ/mole) for
M.sub.2C carbide nucleation contour plot varying at % (Mo) and at %
(V) with superimposed M.sub.2C stoichiometric line at 500.degree.
C. at alloy compositions at % Fe--0.234C--1.32Cu--6.2Ni.
[0029] FIG. 13 is a Mo--V Phase Diagram Section at 900.degree. C.
with alloy composition in atomic %: Fe--0.234C--1.32Cu--6.2Ni. This
diagram shows the phase fields of the FCC austenite and
FCC+V.sub.3C.sub.2 revealing that the M.sub.2C stoichiometric line
is well within the solubility limit.
[0030] FIG. 14 is a graph depicting change in hardness as a
function of alloy copper content for BCC copper strengthening
contribution. Experimental results of other copper strengthened
steels are shown. The dotted line represents the best-fit line for
one-half power law given by equation (5).
[0031] FIG. 15 is a plot depicting room temperature (300K)
austenite stability plotted as a function of Vickers Hardness
Number (VHN). The shaded region shows our range of interest for
austenite stability corresponding to a yield strength requirement
of 150-180 ksi after extrapolation of data from previous alloys,
AF1410 and AerMet100.
[0032] FIG. 16 is a plot of the fraction of Ni in austenite and
phase fraction of austenite in alloy vs. mole fraction of Ni at
500.degree. C. with alloy composition in weight %:
Fe--0.05C--3.65Cu--1.85Cr--0.6Mo--0.1V.
[0033] FIG. 17 is a plot of the equilibrium composition of
austenite as a function of alloy Cr content (wt. fraction) at
510.degree. C.
[0034] FIG. 18 is a quasi-ternary section of the designed
multicomponent alloy system at 510.degree. C. Other alloying
elements are fixed at Fe--0.24C--3.25Cu--6.26Ni--0.35Mo--0.11V (at
%). The tie-triangles shown by thin solid lines indicate
three-phase equilibrium between BCC Cu, austenite and ferrite. The
dashed arrow traces out the trajectory of the austenite phase
composition (solid dots) as a function of increasing alloy Cr
content.
[0035] FIG. 19 is an equilibrium phase fractions at 510.degree. C.
as a function of alloy Cr content (wt fraction).
[0036] FIG. 20 is a plot showing the variation of equilibrium mole
fraction of different phases in the alloy as a function of
temperature, showing that the alloy is solution treatable at
900.degree. C.
[0037] FIG. 21 is a plot of a Scheil simulation for evolution of
the fraction solid with cooling for designed alloy
Fe--0.05C--6.5Ni--3.65Cu--1.84Cr--0.6Mo--0.1V (wt %) in comparison
with equilibrium solidification.
[0038] FIG. 22 is a plot of a Scheil simulation for composition
profile of each alloying element after solidification for designed
alloy Fe--0.05C--6.5Ni--3.65Cu--1.84Cr--0.6Mo--0.1V (wt %). Solid
fraction corresponds to position relative to dendrite arm
center.
[0039] FIG. 23 is a plot of room temperature (300K) stability of
austenite as a function of tempering temperature. The required
stability is predicted for 490.degree. C.
[0040] FIG. 24 is a diagram of the Charpy V-notch impact specimen
dimensions (Standard ASTM E23) with longitudinal axis corresponding
to the L-T orientation.
[0041] FIG. 25 is a diagram of the tensile test specimen dimensions
(Standard ASTM E23).
[0042] FIG. 26 is an optical micrograph of the as-received plate
viewed transverse to the rolling direction at the oxide-metal
interface after etching with 2% natal.
[0043] FIG. 27 is an optical micrograph of the hot-rolled plate
viewed transverse to the rolling direction at the centerline after
etching with 2% natal.
[0044] FIG. 28 is a higher magnification optical micrograph of the
hot-rolled plate at the centerline.
[0045] FIG. 29 is a graph of line profile compositions for
as-received material from oxide-metal interface.
[0046] FIG. 30 is an optical micrograph showing the oxide scale in
the as-received plate.
[0047] FIG. 31 is a plot of relative sample length change and
temperature trace during heating and cooling (quench) cycle from
dilatometry experiment.
[0048] FIG. 32 is a plot of relative sample length change and
temperature trace during heating, cooling and isothermal hold at
377.degree. C. from dilatometry experiment.
[0049] FIG. 33 is a volume fraction evolution of bainite as a
function of time for isothermal temperature of 377.degree. C.
[0050] FIG. 34 is a time-temperature-transformation (TTT) curve for
bainite transformation reaction.
[0051] FIG. 35 is a plot of isochronal (1 hour) tempering response
of prototype alloy. The arrow superimposed on the plot shows that
the design objective is achieved by tempering at 500.degree. C. in
agreement with design prediction.
[0052] FIG. 36 is a plot of isochronal tempering response
represented by Charpy toughness--Vickers hardness trajectory. The
label corresponding to each data point indicates the tempering
temperature.
[0053] FIG. 37 is a plot of Hollomon-Jaffe Parameter correlating
the hardness data obtained for different tempering conditions in
the overaged region.
[0054] FIG. 38 is a SEM micrograph of quasi-cleavage fracture
surface for prototype tempered at 450.degree. C. for 1 hour.
[0055] FIG. 39 is a SEM micrograph of ductile fracture surface for
prototype tempered at 525.degree. C. for 5 hours.
[0056] FIG. 40 is a SEM micrograph of ductile fracture surface
representing toughness enhancement due to transformation toughening
for prototype tempered at 550.degree. C. for 5 hours.
[0057] FIG. 41 is a SEM micrograph of ductile fracture surface
representing toughness enhancement due to transformation toughening
for prototype tempered at 575.degree. C. for 5 hours.
[0058] FIG. 42 is a plot of a multi-step tempering treatments
designed to maximize transformation toughening response represented
by Charpy toughness--Vickers hardness trajectory. The label
corresponding to each data point indicates the tempering time
during the first tempering step. The condition for the second step
is listed on the legend.
[0059] FIG. 43 SEM micrograph of ductile fracture surface
representing toughness enhancement due to transformation toughening
for the 550.degree. C. 30 min+450.degree. C. 5 hrs multi-step
tempering treatment.
[0060] FIG. 44 is a SEM micrograph of a primary void in the
fracture surface of prototype for 550.degree. C. 30 min+450.degree.
C. 5 hrs multi-step tempering treatment.
[0061] FIG. 45 is a plot of true stress--true plastic strain
response. The stress (.sigma.)-plastic strain (.epsilon..sub.p)
behavior is shown by solid lines until uniform elongation and by
dotted line after necking.
[0062] FIG. 46 is a plot of hardness--Yield Strength Correlation
developed from previous data. The heavy black points represent data
from current investigation.
[0063] FIG. 47 is a plot of Charpy impact energy absorbed as a
function of testing temperature for prototype tempered at
550.degree. C. 30 min+450.degree. C. 5 hr. Toughness increment of
30 ft-lb due to dispersed phase transformation toughening is shown.
The toughness band defined by 5 hour and 10 hour single step
tempering is superimposed.
[0064] FIG. 48 is a SEM micrograph of quasicleavage fracture
surface showing flat facets with dimples and tear ridges for the
550.degree. C. 30 min+450.degree. C. 5 hrs multi-step tempering
treatment tested at -84.degree. C.
[0065] FIG. 49 is a SEM micrograph of mixed ductile/brittle mode
fracture surface showing microvoids with some tear ridges for the
550.degree. C. 30 min+450.degree. C. 5 hrs multi-step tempering
treatment tested at -40.degree. C.
[0066] FIG. 50 is a SEM micrograph of purely ductile mode fracture
surface showing primary voids and microvoids for the 550.degree. C.
30 min+450.degree. C. 5 hrs multi-step tempering treatment tested
at -20.degree. C.
[0067] FIG. 51 is a SEM micrograph of purely ductile mode fracture
surface showing primary voids and microvoids for the 550.degree. C.
30 min+450.degree. C. 5 hrs multi-step tempering treatment tested
at 0.degree. C.
[0068] FIG. 52 SEM micrograph of purely ductile mode fracture
surface showing primary voids and microvoids for the 550.degree. C.
30 min+450.degree. C. 5 hrs multi-step tempering treatment tested
at 100.degree. C.
[0069] FIG. 53 is a 3DAP reconstruction for prototype tempered at
450.degree. C. for 1 hour. The elements in the reconstruction are
indicated by their color code. Iron is not shown to provide more
clarity in viewing the particles. z is the direction of
analysis.
[0070] FIG. 54 is a 3DAP reconstruction for prototype tempered at
500.degree. C. 30 min+450.degree. C. 5 hrs. The elements in the
reconstruction are indicated by their color code. Iron is not shown
to provide more clarity in viewing the particles. z is the
direction of analysis.
[0071] FIG. 55 is a 3DAP reconstruction for prototype tempered at
450.degree. C. for 1 hour showing copper precipitates defined at 10
at % isoconcentration surface overlaid on atomic positions of
copper atoms. All other atoms in the reconstruction are not shown.
z is the direction of analysis.
[0072] FIG. 56 is a 3DAP reconstruction for prototype tempered at
500.degree. C. 30 min+450.degree. C. 5 hrs showing copper
precipitates defined by 10 at % isoconcentration surface overlaid
on atomic positions of copper atoms. z is the direction of
analysis.
[0073] FIG. 57 is a proxigram of all the solute species detected in
the 450.degree. C. 1 hr temper specimen with respect to 10 at %
copper isoconcentration surface in the analysis volume.
[0074] FIG. 58 is a proxigram of all the solute species detected in
the 500.degree. C. 30 min+450.degree. C. 5 hrs temper specimen with
respect to 10 at % copper isoconcentration surface in the analysis
volume.
[0075] FIG. 59 is a 3DAP reconstruction for prototype tempered at
500.degree. C. 30 min+450.degree. C. 5 hrs showing austenite
defined by 10 at % Ni level isoconcentration surface overlaid on
atomic positions of nickel and copper atoms. z is the direction of
analysis.
[0076] FIG. 60 is a One-dimensional composition profile along the
atom-probe analysis direction in the 500.degree. C. 30
min+450.degree. C. 5 hrs temper specimen with respect to 10 at %
copper isoconcentration surface in the analysis volume. z is the
direction of analysis.
[0077] FIG. 61 is a toughness-yield strength comparison plot of
Blastalloy160 with other commercial and experimental steels.
DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0078] FIG. 5 presents a schematic process flow of the design
optimization procedure and considerations generally employed to
determine an optimal composition and process for development of
alloys of the invention. Following is discussion regarding these
considerations among others.
Constituent Considerations
[0079] To strengthen the steel while limiting carbon content for
weldability, co-precipitating M.sub.2C carbides and BCC copper have
been employed. By optimizing the particle size and the phase
fraction of the precipitates, the goal of high-strength is
achieved.
[0080] To achieve a goal of 160 ksi yield strength, quantitative
models are employed to relate the contribution from dispersions of
M.sub.2C carbide precipitates and BCC copper precipitates in
secondary-hardened steels. The levels of M.sub.2C carbide formers
and copper are optimized based on the strength contribution from
each of these substructures. Assessment of the yield strength of
the material has been made directly from the hardness data because
of the ease and convenience in measurement of the latter. Hardness
of a material is a direct manifestation of its resistance to
plastic flow, monotonically relating to yield stress. An empirical
relationship has been developed between hardness and yield stress.
The best-fit curve in a log-log plot of hardness vs. yield stress
has been used to determine the relationship based on strain
hardening associated with the alloy. FIG. 6 presents the
experimentally measured hardness--yield stress data superimposed
with the best-fit power-law relationship and the theoretical
straight-line relation describing the same for an ideal plastic
material. The higher hardness of the empirical power-law
relationship relative to the ideal-plastic case represents the
effect of strain hardening, which is more pronounced at lower
strength levels. The point at which the two curves meet represents
the prediction limit of the relationship.
[0081] Thus, the hardness estimate for the target yield strength of
160 ksi from the power-law relationship is 389 VHN. The relation
obtained is: VHN=6.116YS.sup.0.8184 (2) where, VHN (Vickers
Hardness Number) is in kg/mm.sup.2 and YS (Yield Strength) is in
ksi.
[0082] Setting the carbon content of the alloy to ensure good
weldability is appropriate initially prior to evaluation of
hardness factors of the alloy. FIG. 7 presents the Graville diagram
of overall carbon content in the alloy as a function of carbon
equivalent. This shows that at 0.05 wt % C, the steel is not
susceptible to hydrogen--induced cold cracking in the heat affected
zone (HAZ) of weldments. A lower limit C content of about 0.05 wt %
C is desired.
[0083] Based on the predicted change in hardness-carbon content (wt
%) plot shown in FIG. 8, at a C level of 0.05 wt % the hardness
increment due to M.sub.2C carbide precipitation is estimated to be
175 VHN provided a sufficient driving force is maintained to
achieve the particle size range of .about.4-5 nm in FIG. 8. The
base strength of a lath martensitic substructure is estimated as 63
VHN. The additional strength increment of 151 VHN to achieve the
strength goal of 389 VHN is therefore to be attained through BCC
copper precipitation strengthening. The effect of solid solution
strengthening is assumed to be negligible for steels having low
carbon and low hardenability. Thus, the total strength of the alloy
is modeled by breaking it down into three contributions. The
strength is described by the effects of M.sub.2C carbide
precipitates, .tau..sub.M2C; BCC copper precipitates, .tau..sub.Cu
and matrix martensitic structure, .tau..sub..alpha.'.
.tau.=.DELTA..tau..sub.M2C+.DELTA..tau..sub.Cu+.DELTA..tau..sub..alpha.'.-
ident.389VHN (3) The contributions of the individual mechanisms to
achieve the strength goal equivalent to 389 VHN are graphically
presented in FIG. 9. M.sub.2C Carbide Strengthening
[0084] For the high-strength design, it is desired to ensure that
substantially all of the carbon is taken up by the M.sub.2C carbide
formers (Cr, Mo and V) in order to dissolve the cementite in the
matrix since cementite negatively affects strength and toughness.
Therefore, the sum of the atomic concentrations of Cr, Mo, and V is
about double the concentration of C for the M.sub.2C
stoichiometry.
[0085] Compositions are set using the guideline for carbon content
limited to about 0.05 weight % for weldability (FIG. 7); Cu should
be at least 1.5 weight % for significant strengthening, minimum Ni
content should be at least half that of Cu to avoid hot shortness,
and the relative amounts of carbide formers Cr, Mo and V may be
initially set equal in atomic percent. A BCC Cu-rich precipitate is
necessary for Cu precipitation strengthening, an M.sub.2C carbide
phase is necessary for carbide strengthening, and FCC austenite
dispersion is critical for transformation toughening. Referring to
FIG. 10, M.sub.2C solubility is not a limiting factor in the region
of interest, due to the relatively limited C content.
[0086] The stoichiometric constraints of the M.sub.2C carbide
dictate that the total amount of carbide formers (Cr, Mo, V) needed
to balance the carbon content would be 0.468 at %. Using this
constraint, initial plots were constructed of the driving force for
M.sub.2C nucleation vs. at % (Mo) and at % (Cr), setting V at
different levels. FIG. 11 is a representative plot of driving force
contours with varying at % (Mo) and at % (Cr) at an alloy
composition of 0.05 at % V and at 500.degree. C. The stoichiometric
constraint line has been drawn on the plot indicating the line of
allowed compositions for M.sub.2C. Cr has the least effect on
driving force, especially at the higher contents of interest.
[0087] Based on this finding, another set of driving force plots
were created varying at % (Mo) and at % (V) while setting the Cr
level at fixed values. Due to the very small Cr dependence, all the
plots were very similar and so only a representative graph (FIG.
12) was included at 0 at % (Cr). A similar M.sub.2C stoichiometric
line was drawn as before, constraining a maximum driving force at
about 14.4 (kJ/mole). This plot revealed an almost equal effect on
driving force for Mo and V, indicating that any allowed ratio of
the two should give a maximum driving force value, so a series of
calculations were done along the stoichiometric line (maximum
driving force). A feasible alloy composition where all the desired
phases as mentioned before co-exist, is indicated by the dot and
arrow in FIG. 12. The alloy composition in wt % without any Cr is
as follows: Fe--0.05C--1.5Cu--6.5Ni--0.6Mo--0.1V.
[0088] The V--Mo phase diagram section at a solution temperature of
900.degree. C. with the feasible alloy composition was also
calculated. Again, the solubility of these carbide formers is not a
limiting factor in the region of interest. This plot is shown in
FIG. 13.
Copper Precipitation Strengthening
[0089] In addition to M.sub.2C carbide strengthening, BCC copper
precipitation strengthening controls the phase fraction of the
precipitates through the alloy copper content and provides an
additional increment of strength (.ident.151 VHN). The copper
precipitates that contribute to strengthening in steels have a
metastable BCC structure, which are fully coherent with the matrix
having an average diameter of 1-5 nm. The strengthening mechanism
is based on the interaction between the matrix slip dislocation and
the second phase copper-rich particle of lower shear modulus than
the matrix. The shear stress has a maximum value, .tau..sub.max,
given by Equation 3.3. .tau. max = 0.041 .times. Gbf 1 / 2 r 0 ( 4
) ##EQU1## where G is the matrix shear modulus, b is the burgers
vector, f is the volume fraction of atoms and r.sub.0 is the core
radius of the dislocation. Thus the maximum strength that can be
achieved is proportional to the square root of the volume fraction
of the precipitate. Based on this volume fraction dependence of the
precipitate on yield stress, the hardening increment from available
data of copper precipitation strengthened steels was plotted as
shown in FIG. 14. The best-fit line described by a one-half power
law defined the hardening increment dependence on the alloy content
(at %) of copper. .DELTA..tau.(VHN)=83.807X.sub.Cu.sup.1/2 Based on
this relationship, the hardness increment of 151 VHN is achieved by
addition of about 3.25 at % Cu to the alloy composition. Processing
Considerations
[0090] Transformation Toughening
[0091] The toughness of the higher strength steel is improved by
utilizing the beneficial properties of Ni-stabilized precipitated
austenite. This form of austenite can precipitate during annealing
or tempering at elevated temperatures above about 470.degree. C.
The fact that this dispersed austenite forms by precipitation is
significant because it allows greater overall control of the amount
and stability of the austenite. Further processing and treatments
can be used in the form of multi-step tempering to first nucleate
particles in a fine form at a higher tempering temperature and then
complete Ni enrichment during completion of precipitation
strengthening (cementite conversion to M.sub.2C) at a lower final
tempering temperature.
[0092] The austenite dispersion has stability and formation
kinetics to ensure maximum toughening enhancement. Other factors
controlling the stability of austenite are particle size and stress
state sensitivity, the latter being related to the transformational
volume change. Stability of an austenite precipitate is defined by
chemical and mechanical driving force terms. At the
M.sub.s.sup..sigma. temperature (where transformation occurs at
yield stress) for the crack-tip stress state, the total driving
force equals the critical driving force for martensite nucleation,
as represented by Equation 6. .DELTA. .times. .times. G ch +
.sigma. y .times. .times. d .DELTA. .times. .times. G .sigma. d
.sigma. .times. | cracktip = - [ 2 .times. .times. .gamma. nd + G 0
+ W f ] ( 6 ) ##EQU2## Rearranging the terms and substituting a
dependence of defect potency on particle volume V.sub.p, defines a
convenient stability parameter: .DELTA. .times. .times. G ch + W f
+ K ln .function. ( V p ) = - [ .sigma. y .times. d .DELTA. .times.
.times. G .sigma. d .sigma. .times. | cracktip .times. + G 0 ] ( 7
) ##EQU3## .DELTA.G.sup.ch is the transformation chemical free
energy change and W.sub.f is the athermal frictional work term
described in Section 2.4. .DELTA.G.sup.ch is temperature and
composition dependent while W.sub.f is only composition dependent.
W.sub.f will vary with tempering temperature due to the change in
austenite composition. .sigma..sub.y is the yield stress of the
material, .DELTA.G.sup.ch is set by the stress state and G.sub.0 is
a nucleus elastic strain energy term. K is a proportionality
constant, .gamma. is the nucleus specific interfacial energy and d
is the crystal interplanar spacing.
[0093] The austenite stability for a given set of conditions or
service temperature for a given dispersion can be assessed by the
parameter given by the left-hand side of Equation 7. Austenite
stability parameter becomes the sum of the chemical driving force
for transformation of FCC austenite to BCC martensite at room
temperature (300K) and the frictional work term for martensitic
interfacial motion: .DELTA.G.sup.ch+W.sub.f. The model is
represented in Equation 8. W f = i .times. ( K i .times. X i 1 / 2
) + j .times. ( K j .times. X j 1 / 2 ) + k .times. ( K k .times. X
k 1 / 2 ) + K Co .times. X Co 1 / 2 ( 8 ) ##EQU4## where the K's
represent the coefficients used to fit the solid solution
strengthening data and i=C, N; j=Cr, Mn, Mo, Nb, Si, Ti, V; and
k=Al, Cu, Ni, W. Equation 8 further indicates that the stability
parameter is a linear function of the yield strength of the
material.
[0094] FIG. 15 gives the plot of the austenite stability parameter,
.DELTA.G.sup.ch+W.sub.f, at room temperature against Vickers
hardness of the alloy. The room temperature stability of the
austenite dispersion projected from the hardness (or strength)
requirement of the design is marked by the shaded region in the
figure and quantitatively expressed in Table 1. To achieve a goal
of 160 ksi yield strength equivalent to Vickers hardness of 389
(Rc40 equivalent), the estimated optimum .DELTA.G.sup.ch+W.sub.f
value of 2837 J/mole is found for the required stability.
TABLE-US-00001 TABLE 1 Target Chemical Driving Force
(.DELTA.G.sup.ch) + Frictional Work (W.sub.f) Value Rockwell C
Hardness Vickers Hardness .quadrature.G.sub.ch + W.sub.f Alloy
R.sub.c VHN (kg/mm.sup.2) J/mol AerMet 100 54 577 4350 AF1410 48
484 3600 Design 40 389 2837
[0095] Plots of both the phase fraction of austenite and nickel
content in the austenite phase vs. alloy atomic fraction Ni were
computed. FIG. 16 was calculated at an estimated final tempering
temperature of 500.degree. C. for substitutional diffusion and
revealed that a minimum of 3.5 at % Ni is required to get austenite
and a maximum fraction of nickel in the austenite of about 0.30
could be obtained. It also showed that at the 6.25 at % Ni
composition, about a 0.10 phase fraction of austenite would be
formed as shown by the arrows. Thus, the alloy Ni level is set to
about 6.25 at %, which also saturates the austenite Ni content to
30 at %.
Alloy Composition And Processing Integrated
[0096] The overall composition was optimized so that all of the
phases necessary for strengthening and toughening are
simultaneously present. The maximum M.sub.2C driving force is
obtained with no chromium. The copper added for precipitation
strengthening went instead into the austenite phase. A study of the
equilibrium austenite composition with varying alloy Cr content was
then undertaken as given in FIG. 17. It was found that Cr
partitions Cu out of austenite and into the BCC precipitate phase
effectively at 2 wt % and above.
[0097] To understand the effect of Cr on partitioning of Cu out of
the austenite phase, a detailed investigation was done based on a
quasi-ternary section of the multicomponent system at 510.degree.
C. as presented in FIG. 18. The equilibrium Cr and Cu phase
compositions of BCC Cu, austenite and the ferrite phases connected
with tie-triangles are presented for different Cr contents of the
alloy. The austenite equilibrium point abruptly shifts to a much
lower Cu level with an increase of alloy Cr level from 1 at % to
1.2 at %. This confirms that Cu partitions to the BCC precipitate
phase above 1.2 at % Cr. Thus, 2 at % Cr (equivalent to 1.84 wt %
Cr) was set for the alloy to make the copper in the alloy available
for precipitation strengthening.
[0098] The relative fractions of the different phases in the
microstructure were then calculated as a function of the alloy Cr
content to confirm the effect of chromium as shown in FIG. 19. This
confirms that at 2 wt % Cr, there is sufficient precipitation of
austenite (>0.1 mole fraction) for transformation toughening and
bcc Cu (.about.0.03 mole fraction) for strengthening.
Processing Factors
[0099] Solution Treatment Temperature and Allotropic
Transformations
[0100] A solution treatment temperature of 900.degree. C. was
chosen. With the increased levels of Cu and Cr it was confirmed
that the alloy was solution treatable at 900.degree. C. as shown by
the phase fraction plot in FIG. 20.
[0101] For this alloy composition, the martensite and bainite
kinetic models predict an M.sub.s temperature of 298.degree. C. and
a bainite start (B.sub.s) temperature of 336.degree. C. These are
deemed sufficiently high to ensure formation of bainite/martensite
mixtures with air-cooling.
Microsegregation Behavior
[0102] Solidification of alloys generally occurs with segregation,
which can have a strong effect on the alloy's final properties.
[0103] FIG. 21 presents the solidification simulation result as
temperature vs. fraction solid using a non-equilibrium Scheil
simulation and compares it with the full equilibrium case. FIG. 22
presents the composition profiles calculated by a Scheil simulation
showing the degree of microsegregation in the solid after
solidification. Here, the fraction of solid is equivalent to a
position relative to a dendrite arm center. The results presented
in Table 2 predict that Mo has the greatest potential for
segregation. However, since the level of Mo in the alloy is low, no
serious microsegregation problems are encountered. TABLE-US-00002
TABLE 2 Amplitude of microsegregation with respect to each alloying
element predicted by Scheil simulation at 95% solidification
Alloying Elements Ni Cu Cr Mo V Nominal Alloy Composition - 6.38
3.31 2.04 0.36 0.11 C.sub.alloy (at %) Microsegregation 1.29 1.67
0.72 0.34 0.05 Amplitude C.sub.0.95-C.sub.0 (at %)
Tempering Temperature
[0104] The austenite stability for this transformation toughened
alloy is dependent on the optimal tempering temperature condition.
With the alloy composition fixed, the austenite stability is
calculated as a function of tempering temperature as shown in FIG.
23. It illustrates that the .DELTA.G.sub.ch+W.sub.f value of 2836
J/mole desired for this alloy is achieved for a tempering
temperature of 490.degree. C., very close to the originally assumed
temperature of 500.degree. C.
[0105] A composition is thus in a preferred embodiment for the
ultratough, high strength weldable plate steel (in wt %) to be
tempered at 490.degree. C. of about:
[0106] Fe--0.05C--3.65Cu--6.5Ni--1.84Cr--0.6Mo--0.1V.
[0107] The composition should be solution treatable at 900.degree.
C., with predicted Ms and Bs transformation temperatures of
298.degree. C. and 336.degree. C. respectively. Initial tempering
at a slightly elevated temperature will help nucleate the austenite
before tempering at 490.degree. C. to enrich the Ni content to the
designed level.
Experimental Results and Examples
[0108] Material
[0109] Special Metals Corporation in New Hartford, N.Y. produced
the alloy in a 34-pound heat by Vacuum Induction Melting (VIM) from
100% virgin raw materials and cast into
9.5''.times.8''.times.1.75'' (24.1 cm.times.20.3 cm.times.4.5 cm)
slab ingots as a simulation of a continuous casting process. The
as-cast ingot was subsequently homogenized at 2200.degree. F.
(1204.degree. C.) for 8 hours and then hot rolled to 0.45'' (1.1
cm) thickness followed by air-cooling to room temperature by
Huntington Alloys in Huntington, W. Va. The final dimension of the
plate measured roughly 33''.times.10''.times.0.45'' (83.8
cm.times.25.4 cm.times.1.1 cm). The hot-rolled plate was annealed
at 900.degree. F. (482.degree. C.) for 10 hours to improve
machinability of the material. The designed and the actual
compositions (in wt %) of the alloy is given in Table 3. The
impurity level in the alloy was measured as 0.002 wt % S, 13 ppm O
and 2 ppm N. TABLE-US-00003 TABLE 3 Designed and Measured
Composition (in wt. %) of alloy Alloy Fe C Cu Ni Cr Mo V Design
Bal. 0.05 .+-. 0.01 3.65 .+-. 0.05 6.5 .+-. 0.2 1.84 .+-. 0.05 0.6
.+-. 0.05 0.1 .+-. 0.01 Measured Bal. 0.040 3.64 6.61 1.78 0.58
0.11
[0110] Experimental Procedures [0111] Heat Treating
[0112] All samples were solution treated at 900.degree. C. for 1
hour and quenched in water followed by a liquid nitrogen cool for
30 minutes prior to every tempering treatment to ensure a fully
martensitic starting microstructure and eliminate any retained
austenite. Solution treatments were done in an argon atmosphere to
prevent oxidation of samples. To ensure rapid heating of the entire
sample, the short-time nucleation stage heat treatments were
conducted using a molten salt bath followed by water-quenching to
room temperature. The salt used for the molten bath was
Thermo-Quench Salt (300-1100.degree. F.) produced by Heat Bath
Corporation. The residue layer from the salt pot treatment was
ground off before the second step aging treatment. The standard
aging treatments for longer times (1-10 hours) were performed in a
box furnace under vacuum (to prevent oxidation and decarburization)
and then air-cooled to room temperature. Vacuum was achieved by
encapsulating the samples in 0.75'' diameter pyrex tubes connected
to a vacuum system. The pyrex tubes were evacuated by a mechanical
roughing pump followed by a diffusion pump. During evacuation, the
tubes were backfilled with argon three times before reaching a
final vacuum of <5 mtorr. Each tube was then sealed with an
oxygen/propane torch. [0113] Metallographic Sample Preparation
[0114] All samples were ground and polished directly to 1 .mu.m
finish using a Buehler Ecomet-4 variable speed automatic
grinder/polisher. The samples prepared for measuring hardness were
mounted in room temperature curing acrylic, while those prepared
for microsegregation studies were hot mounted with conductive
phenolic resin using a Stuers LaboPress-1 after nickel-plating for
edge retention of the oxide layer during grinding and polishing.
Microsegregation samples were etched by submersion in a 2% nital
(2% nitric acid in ethanol) solution for 10-30 seconds to reveal
the compositional banding close to the metal-oxide interface
associated with scale formation during hot working. Following
etching, the samples were viewed with an optical microscope to
study the banded structure in the as-cast material. [0115]
Dilatometry
[0116] Dilatometry is used to study phase transformations by
recording length changes versus temperature. For these studies a
computer controlled MMC Quenching Dilatometer was used. Specimens
were prepared by EDM (Electro-Discharge Machining) wire cutting
into cylindrical rods 10 mm long and 3 mm in diameter. The samples
are heated by an induction furnace and cooled by jets of helium
gas. They are mounted between two low expansion quartz platens,
which are lightly spring-loaded and are connected to an LVDT
transducer that records the length. The temperature is monitored by
a Pt-Pt 10% Rh thermocouple spot welded directly to the sample
surface. The sample stage is enclosed in a vacuum chamber connected
to a turbo-mechanical pump and mechanical backing pump capable of
achieving a vacuum of 10.sup.-4 torr.
[0117] Dilatometry was used for determining the martensite start
temperature (M.sub.s) and for evaluating the bainite transformation
kinetics. For estimating the experimental M.sub.s temperature,
samples were heated at a rate of 2-3.degree. C./sec to 1050.degree.
C., held for 5 minutes for homogenization and then rapidly quenched
(>100.degree. C./sec) to room temperature. The M.sub.s
temperature was determined as the transition at which the sample
started expanding on cooling. For studying the bainite kinetics,
samples were held isothermally for 2 hours at bainite
transformation temperatures between 360-420.degree. C. after
quenching (Cooling rate from 800.degree. C. to 500.degree. C.,
T.sub.8/5=50.degree. C./sec) from the austenizing temperature. The
length change at the isothermal hold temperature is a measure of
the amount of bainitic transformation. All samples were austenized
at 1050.degree. C. for 5 minutes and then rapidly quenched prior to
the actual runs of martensite and bainite transformation in order
to ensure uniform starting microstructure. [0118] Microhardness
Testing
[0119] Vickers hardness was measured for every aging condition as a
measure of strength. The relationship between hardness and yield
strength helped to assess the mechanical properties directly from
the hardness data. Hardness measurements of materials in this study
were performed using the Buehler Micromet II Micro Hardness Tester
based on the method prescribed in ASTM standard E384. A diamond
Vickers pyramidal indenter with face angles of 136.degree. is used
to make the indentations. After applying a load of 200 g for 5
seconds, the diagonals of the indent were measured at 40033
magnification to obtain the Vickers Hardness (VHN) according to
Equation 9. VHN = 1.854 .times. P d 2 ( 9 ) ##EQU5## where P is the
load in kg. and d is the average length of the diagonal in
millimeters of the indent. Prior to testing, all the heat-treated
samples were mounted in acrylic mold and polished to 1 .mu.m. The
samples were at least 8 mm thick and ground to reveal opposite
surfaces to avoid any errors due to anvil effects. At least ten
hardness measurements were recorded uniformly across the
cross-section for every sample tested and the average was
documented as the hardness value. [0120] Impact Toughness
Testing
[0121] The impact toughness properties for the different heat
treatment conditions of the alloy were measured using a Tinius
Olsen 260 ft-lb (352 J) impact-testing machine. Prior to testing,
the samples were machined according to the ASTM standard Charpy
V-notch dimensions (1996 ASTM E23) 10 mm.times.10 mm.times.55 mm
(0.39''.times.0.39''.times.2.17'') with a 45.degree. notch of depth
2 mm and root radius of 0.25 mm placed at the center of the long
side. The longitudinal axis of the specimen corresponded to the L-T
orientation. A schematic view of the sample geometry is given in
FIG. 24. The impact fracture energy was measured directly on analog
scale and the given impact energy data was mostly based on a
two-sample average. Most impact properties were evaluated at room
temperature. For the low temperature impact fracture properties,
the aged samples were held for 20 minutes at the test temperature
in an Instron low temperature furnace connected to a liquid
nitrogen supply. Within 5 seconds of removal from the furnace, the
samples were placed inside the machine and struck with the 100-lbf
hammer. [0122] Tensile Testing
[0123] Tensile test specimens were machined from blanks measuring
approximately 10 mm.times.10 mm.times.70 mm
(0.39''.times.0.39''.times.2.76'') from the original plate parallel
to the longitudinal rolling direction. Prior to machining, the
samples were solution-treated and aged as discussed in Section
2.2.1. From each blank, sub-sized tensile specimens, scaled in
accordance to ASTM standards (1996 ASTM E8M) were machined as shown
schematically in FIG. 25. The final specimen had a gage diameter of
6 mm (0.24'') and a gage length of 30 mm (1.18'').
[0124] All tensile tests were performed at room temperature using a
computer controlled Sintech 20/G screw driven mechanical testing
machine with a 20,000 lb (8896 N) load cell at a constant
cross-head speed of 0.005 in/sec (0.127 mm/sec). The load cell was
calibrated prior to every data set using the computer controlled
calibration test. A calibrated extensometer of gage length 1''
(25.4 mm) was attached to the sample during testing to measure the
displacement. The load-time response was recorded using the
TestWorks computer software package interfaced with the Sintech
tensile testing machine. The actual cross-sectional areas and gage
lengths of the specimens were measured prior to testing and listed
in the testing program. Area reduction and extension were measured
manually upon completion of the test. Engineering stress-strain
curves were obtained directly thorough the TestWorks program. Based
on a two-sample average for select processing conditions, the
ultimate tensile strengths, 0.2% offset yield strengths and total
elongations were obtained. [0125] Scanning Electron Microscopy
[0126] A Hitachi S-3500 scanning electron microscope (SEM) with a
tungsten hairpin filament was used to investigate the composition
banding in the as-rolled samples and the fracture surfaces of the
Charpy impact specimens. The microscope uses Quartz PCI Image
Management Software for capturing images and for conducting
quantitative analysis. For analysis, the samples were mounted on
graphite tape and examined in the SEM with a 20 kV electron beam at
a vacuum level of 10.sup.-4 torr inside the specimen chamber. The
secondary electron (SE) detector was used for imaging both the
etched and fracture surfaces. The compositionally banded structure
of the etched sample was characterized quantitatively from the
metal-oxide interface using the PGT Energy Dispersive X-ray
analyzer with digital pulse processing. Fractography analysis was
done to characterize the fracture surface and micrographs
containing interesting features were taken. [0127] Atom Probe/Field
Ion Microscopy (AP-FIM)
[0128] A three-dimensional atom probe microscope was used for
characterizing the size, number-density and composition of
nanoscale strengthening (Cu precipitates) and toughening
(Ni-stabilized austenite) dispersions in the heat-treated samples.
The atom probe, operated and maintained under an ultra-high vacuum
system (10.sup.-10-10.sup.-11 torr) combined with a field ion
microscope, operated with imaging gas at a pressure level of
10.sup.-5 torr, makes it an extremely high-resolution microscopy
technique.
[0129] The specimens (atom probe tips) were prepared by a two-step
electropolishing sequence of small rods (100 mm long with 200
.mu.m.times.200 .mu.m square cross-section) cut from heat-treated
hardness samples. Initial polishing was done using a solution of
10% perchloric acid in butoxyethanol at room temperature applying a
DC voltage of 23V until the square rods were shaped into long
needles with a small taper angle. A solution of 2% perchloric acid
in butoxyethanol at room temperature was used for necking and final
polishing to produce a sharply pointed tip, with a radius of
curvature less than 50 nm. The voltage was gradually decreased from
12V DC to 5V DC during the final stages of electropolishing.
[0130] Each atom probe specimen of tip radius 10 to 50 nm is raised
to a high positive potential of 5-15 kV, resulting in an
exceptionally strong electric field on the order of 50 V/nm. FIM
analysis was performed at temperatures between 50K-80K with a
chamber pressure of 10.sup.-5 torr consisting of neon gas. The
voltage on the tip was raised until an FIM image was observed on
the monitor. Neon atoms, which are used as an imaging gas for
steel, are ionized in the high electric field causing the
positively charged ions to accelerate to a microchannel plate
array. The ionization process occurs at prominent atomic sites at
the edge of a crystallographic plane corresponding to a particular
atom. A continuous stream of ions forms an image on a phosphorus
screen that represents the nanometer-scale structure of the
specimen tip. FIM images were captured during analysis using the
Scion Image imaging software. For atom probe analysis, the specimen
is then rotated towards the reflectron for aligning the primary
detector on the region of interest in the FIM image (usually near a
pole or on a precipitate in the FIM image). Atom probe analysis is
then conducted at temperatures 50K and 70K under ultra-high vacuum
conditions (10.sup.-10-10.sup.-11 torr) for pulsed
field-evaporation with a pulse fraction (pulse voltage/ steady
state DC voltage) of 20% at a pulse frequency of 1500 Hz.
[0131] Atom probe microanalysis is the study of the specimen
composition by pulsed evaporation. Field evaporation of the
specimen occurs at higher electric fields than ionization of
imaging gas ions. The positively charged ions evaporated from the
specimen are accelerated towards a detector. By measuring the time
of flight, it is possible to determine the mass to charge ratio of
the ions according to the following equation: m n = k .function. (
.alpha. .times. .times. V dc + .beta. .times. .times. V pulse )
.times. ( t + t 0 d ) 2 ( 10 ) ##EQU6## where m is the atomic mass,
n is the charge, k is a constant related to the elementary charge
of an electron, V is the DC or pulse voltage, t is the time,
t.sub.0 is a time offset from electronic delays, and .alpha. and
.beta. are system specific calibration parameters.
[0132] The standard error, .sigma., for compositions measured using
an atom probe is calculated using binomial statistics to account
for the statistical uncertainty associated with small sampling
sizes according to the equation: .sigma. = c i .function. ( 1 - c i
) N ( 11 ) ##EQU7## where c.sub.i is the measured composition of
element i and N is the total number of ions sampled. This standard
error does not account for any overlapping mass to charge ratios
between different elements. Systematic errors that may interfere
with the collection of specific elements such as carbon may be an
additional source of error.
[0133] Three-dimensional atom probe (3DAP) records the
two-dimensional location of atoms and determines the third
dimension (z) by the sequence of arrival of atoms to the detector,
thus providing a three-dimensional reconstruction of the specimen
tip. The evaporated ion collides with a primary detector that
records the time of flight, and the phosphorus screen emits light.
The light is split by a partially silvered mirror at 45.degree. to
both a camera and an 8 by 10 array of anodes which determine the
position of the ion.
[0134] The data from 3DAP was analyzed and visualized by the
software ADAM developed by Hellman et al. Different elemental
isotopes were distinguished by their mass/charge ratio. The overlap
of isotope masses between elements contributed to the experimental
error in addition to the statistical counting error. A range of
tools is available in ADAM to analyze the data from the 3DAP. One
feature of ADAM is the ability to define planar or cylindrical
regions of interest and to perform analyses such as concentration
profiles, ladder diagrams and composition maps with respect to that
region of interest. For the data containing copper precipitates,
varying in composition from the matrix, it was possible to define
isoconcentration surfaces of constant composition. The
three-dimensional representation of these isoconcentration surfaces
allows for a qualitative view of the approximate size and shape of
the precipitates being studied. ADAM has been designed to employ
this method by creating a discrete lattice of nodes for which the
local composition is calculated. The isoconcentration surfaces then
have discrete positions. The creation of isoconcentration surfaces
allows for another method of 3DAP data analysis referred to as the
proximity histogram, or proxigram. The minimum distance to an
isoconcentration surface is calculated for each ion in the data set
and the ions are then assigned to bins according to distance. The
concentration of each bin is calculated and plotted as a function
of distance to the isoconcentration surface. The standard error of
each bin is calculated and displayed on the proxigram.
Experimental Evaluation
[0135] The analysis began with evaluation of the processability
characteristics of the designed alloy at an experimental-heat
scale. Optimization of the tempering response of the alloy designed
for multi-step treatment helped to attain a significant
toughness/strength combination characterization of the
strengthening and toughening dispersions related the structure to
the properties.
[0136] Primary Processing Behavior [0137] Microsegregation and
Hot-working Behavior
[0138] The achievement of the property objectives begins with
meeting the initial processability requirements, i.e., castability
of the steel. Microsegregation is a common problem observed in
high-alloyed castings and hot-worked products, which limits the
mechanical properties.
[0139] To study the microsegregation behavior in the cast alloy,
the as-received material (homogenized for 8 hours at 1204.degree.
C., hot-rolled for 75% reduction to 0.45'' or 4.5 cm thick plate
and then annealed at 482.degree. C. for 10 hours) in the form of a
10 mm.times.10 mm.times.20 mm sample, was etched with 2% nital
following standard metallographic polishing to 1 .mu.m. Low
magnification transverse optical micrographs revealed both the
banded structure oriented along the longitudinal rolling direction
and the oxide-metal interface as shown in FIG. 26.
[0140] The centerline of the hot-rolled plate did not reveal as
much of a banded structure as the surface region, as shown in FIG.
27. Higher magnification optical micrograph at the centerline of
the plate presented in FIG. 28 shows an equiaxed microstructure,
which is predominantly lath martensite in the form of packets
within the prior austenite grain boundaries of an average size of
.about.50 .mu.m.
[0141] The composition bands revealed on etching in FIG. 26 were
estimated to be of 40-50 .mu.m thickness. The extent of
microsegregation within these bands was determined by measuring the
composition profile across the thickness of the plate near the
oxide-metal interface. Composition data was collected every 4 .mu.m
starting from the metal-oxide interface and proceeding towards the
center of the plate. The composition variation across the bands
with respect to the major alloying elements Ni, Cu, Cr and Mo is
presented in FIG. 29. It was found that compositional banding in
the plate was limited to an amplitude of approximately 6-7.5 wt %
Ni, 3.5-5 wt % Cu, 1.6-2 wt % Cr, and 0.2-0.5 wt % Mo. From the
strength model, a variation in the level of Cu across the bands
within 3.5 to 5 wt % corresponds to a predicted hardness variation
of 30 VHN equivalent to 6.8 ksi (.about.47 MPa) in yield strength.
This will promote a smooth yielding behavior as confirmed by the
tensile property behavior.
[0142] Another important factor determining the processability of
an alloy is the material response during high temperature
deformation or formability. Hot shortness is a common problem
associated with high copper steel production. During the rolling
stage of the fabrication process, the effect of hot shortness is
observed by the appearance of surface cracks or fissures leading to
unacceptable products. At hot rolling temperatures above
1050.degree. C. in an oxidizing atmosphere, iron is selectively
oxidized leaving an enrichment of copper near the oxide-metal
interface. If the composition of the copper enriched region exceeds
the liquid-austenite equilibrium limit, the copper enriched liquid
phase enters the grain boundary of the austenite causing
intergranular fracture during hot rolling. A high Ni/Cu ratio of
1.8 was maintained to prevent any hot-shortness problems during
processing. Successful hot rolling of the alloy was demonstrated
during processing. As further verification, the oxide layer of the
as-received material was examined carefully for any evidence of
Cu-rich regions. FIG. 30 shows an optical micrograph of the oxide
layer in the as-received plate. The oxide-metal interface does not
show any evidence of hot shortness. Composition analysis of various
regions in the oxide layer did not reveal any Cu rich phase but did
show some Ni-enriched phases varying from 20 to 80% within the
Fe-rich oxide. This study thus supports the ability of Ni to cause
occlusion of the Cu-enriched liquid during oxidation. [0143]
Evaluation of Allotropic Kinetics
[0144] A dilatometry study was conducted to determine the
allotropic kinetics of the prototype. The first step involved the
measurement of the martensite start temperature (M.sub.s) of the
designed alloy. FIG. 31 presents a plot of the relative length
change vs. temperature, used to determine the transformation points
during the heating and cooling (quench) cycle of a dilatometry
experiment. Straight lines are fit to the single phase portions of
heating and cooling curves, the full width between them defining
full transformation. The series of dashed lines superimposed on the
length and temperature trace represent varying degrees of partial
martensitic transformation during rapid quench from an austenizing
temperature of 1050.degree. C. The threshold for transformation is
taken as 1%. Thus, M.sub.s was determined from the 1% martensitic
transformation point as shown in FIG. 31. The M.sub.s temperature,
averaged over 15 dilatometry runs, is 360.+-.8.4.degree. C.
[0145] Since the alloy is a bainite/martensite microstructure
during air-cooling of plates, the bainite kinetics was determined
by studying the isothermal time-temperature-transformation
characteristics of the steel through dilatometry. This information
is useful in determining the processing necessary in order to
achieve bainitic transformation of 50%, for example. The amount of
bainitic transformation was determined by isothermal hold
experiments (after an initial quench step) performed at incremental
temperatures above the martensite start temperature. This data was
then compiled and analyzed in order to plot a
time-temperature-transformation (TTT) curve.
[0146] The relative length change vs. temperature dilatometry trace
for a two-hour isothermal hold at 377.degree. C. is presented in
FIG. 32. The percent of bainitic transformation is determined by
measuring the length increase upon arrival at the isothermal hold
temperature and dividing it by the total FCC(.gamma.)-BCC(.alpha.)
length difference (defined from the martensitic transformation in
FIG. 31) at the isothermal hold temperature. In this case, the
total bainitic transformation that took place after 2 hours is
44.1%. The evolution of bainitic transformation with respect to
time can be determined from the measurement presented in FIG. 32.
This behavior at 377.degree. C. is shown in FIG. 33. From this plot
it is apparent that the volume fraction of bainite is saturated
after a two-hour isothermal hold. Similar analyses were carried out
for each two-hour test performed at isothermal temperature ranging
from 362.degree. C. to 407.degree. C. Table 4 summarizes the
maximum transformation levels at all temperatures. The TTT curve
was then determined by analysing the data at each isothermal hold
temperature. For example, a 1% transformation curve is plotted by
finding the time at which the sample exhibits 1% transformation at
different isothermal temperatures. The TTT curve based on the data
from all of the isothermal runs is presented in FIG. 34. It shows
achievement of a 50% bainite/martensite mix in approximately 4
minutes at 360.degree. C. The experimental B.sub.s temperature was
determined to be 410.degree. C., 50.degree. C. higher than the
corresponding M.sub.s temperature (360.degree. C). TABLE-US-00004
TABLE 4 Saturation volume fraction of bainite as a function of
isothermal temperature Temperature (C.) Saturation Volume Fraction
of Bainite 362 0.609629 367 0.526697 372 0.5003 377 0.440938 382
0.242981 387 0.265119 392 0.098382 402 0.015628 407 0.007966
[0147] Thermal Process Optimization [0148] Isochronal Tempering
Response
[0149] An isochronal tempering study was conducted to evaluate the
tempering characteristics of the alloy and provide a baseline for
multi-step tempering treatments. For simplicity, the tempering
response investigation was done in a uniform martensite matrix to
minimize retained austenite effects. Deleterious transformation
products from retained austenite decomposition during tempering
could negatively affect the toughness. After a solution treatment
at 900.degree. C. for 1 hour followed by a water quench and liquid
nitrogen cool, tempering was performed for 1, 5 and 10 hours under
vacuum. Samples were finish machined, notched and then tested at
room temperature for Charpy impact toughness. Hardness measurements
were taken directly from the polished surface of the Charpy
specimens.
[0150] The tempering response for 1 hour isochronal tempering was
investigated over a temperature range of 200.degree. C.-600.degree.
C. in the solution-treated alloy and is shown in FIG. 35. The
1-hour isochronal tempering study demonstrates that a peak hardness
level is reached at 420.degree. C. followed by gradual overaging.
The retention of high hardness even after the peak aging condition
to 500.degree. C. can be attributed to precipitation of M.sub.2C
carbides and a fine austenite dispersion. The hardness at
500.degree. C. is (represented by an arrow in FIG. 14).
[0151] After confirming the basic secondary hardening
characteristics of the alloy, a series of isochronal tempering
treatments of Charpy specimens were done for 1, 5 and 10 hours
within a temperature range of 400-600.degree. C. FIG. 36
illustrates the room temperature Charpy toughness
(C.sub.v)--Vickers hardness (VHN) trajectory for the indicated
tempering temperatures. This establishes the baseline of the
toughness-hardness (strength) combination in tempered martensitic
microstructures.
[0152] At the shortest tempering time of 1 hour, FIG. 36
demonstrates that cementite formation limits toughness, and as Cu
precipitates in its presence, strength increases from 400.degree.
C. to 450.degree. C. tempering treatment while there is a sharp
decline in toughness. With further tempering, cementite begins to
dissolve as a result of M.sub.2C carbide formation in combination
with BCC copper precipitation at the peak aging condition. This
results in an increase of both strength and toughness. The
toughness-hardness trajectory takes a sharp turn thereafter, as the
strengthening precipitates begin to coarsen exceeding their optimum
sizes and the strength continues to decrease with averaging. FIG.
36 suggests that peak hardness occurs at 450.degree. C. 5 hour
tempering and the corresponding toughness resides on an upper band
indicating complete dissolution of paraequilibrium cementite by
precipitation of an optimal size M.sub.2C strengthening
dispersion.
[0153] The highly overaged region is also likely associated with
precipitation of a fine dispersion of austenite, which increases in
stability due to Ni enrichment at higher tempering times. A feature
observed in the toughness-hardness trajectory for 5 hour tempering
in FIG. 36 between tempering temperatures of 525.degree. C. and
575.degree. C. is a toughness enhancement from the baseline
toughness of 144 ft-lbs to 170 ft-lbs respectively, a toughness
increment by 18% at a strength level corresponding to 355 VHN. This
is characteristic of the transformation toughening phenomenon
caused by the austenite reaching an optimal stability for the lower
strength condition.
[0154] The tempering response of the hardness (strength) can be
correlated to an empirical Larson-Miller type parameter, known as
the Hollomon-Jaffe tempering parameter. The parameter is defined as
T(18+1n(t)) where T is tempering temperature in K and t is the
tempering time in minutes, and is used for correlation of hardness
data at higher tempering temperatures between 400.degree. C. and
600.degree. C. FIG. 37 presents the measured values of hardness for
different tempering conditions as a function of the Hollomon-Jaffe
tempering parameter. Fairly good agreement with the parameter is
obtained for hardnesses under overaged tempering conditions. The
parameter can provide a simple interpolation scheme to adjust
tempering for a desired strength level.
[0155] The fracture surfaces of the broken Charpy impact testing
samples were observed under SEM to characterize the mode of
fracture. The fracture surface for the 450.degree. C. 1 hour
tempering condition is presented in FIG. 38. The SEM micrograph
reveals that the sample failed by quasi-cleavage fracture with
signs of intergranular embrittlement. Quasi-cleavage is
characterized by an array of cleavage failures connected by ductile
tear ridges but is a much more desired fracture mode compared to
intergranular fracture. The fracture mode represents relatively
brittle behavior attributed to the presence of undissolved
cementite at short tempering times.
[0156] For higher tempering times and temperatures, ductile
fracture occurred by microvoid nucleation and coalescence.
Representative SEM micrographs showing ductile mode of fracture for
5 hour tempering marked by toughness enhancement due to
transformation toughening in FIG. 36 are presented in order of
increasing tempering temperature in FIGS. 39 through 41. FIG. 39
clearly shows that a completely ductile mode of fracture is
achieved with 5 hour 525.degree. C. tempering and micrographs
presented in FIGS. 40 and 41 represent fracture surfaces with
increased toughness due to transformation toughening, indicated by
the relatively higher degree of primary void growth. [0157]
Toughness Optimization by Multi-step Tempering
[0158] Heat treatment for stabilization of austenite for dispersed
phase transformation toughening phenomenon is directed towards
combined size refinement and compositional enrichment of the
austenite particles. A two-step tempering process consisting of an
initial high temperature, short time treatment followed by an
isothermal tempering treatment is employed to achieve this goal.
The first step is designed to nucleate a fine, uniform dispersion
of intralath austenite and strengthening particles of sub-optimal
size formed directly by increasing the driving force for
precipitation. This is achieved by a short time, high-temperature
tempering step designed to give an underaged state based on the
isochronal tempering study. At this stage, it is advised to
understand the implications of the kinetic competition between the
precipitation of austenite and strengthening dispersions namely,
BCC copper and M.sub.2C carbides. In the alloy, the austenite
precipitation kinetics is slower than the BCC copper precipitation
kinetics, which in turn is considerably slower than the carbide
precipitation process at intermediate tempering temperatures. It
is, therefore, desired to optimize the time for the
high-temperature austenite nucleation step, since the carbides
might become overaged at higher times and full hardness cannot
likely be achieved. Yet this uncertainty in loss of strength by
averaging of carbides is overcome in the alloy because of
additional strengthening of nearly 40% provided by BCC copper
precipitation, which has slower coarsening kinetics than the
carbides. The second tempering step is thus optimized to enhance
Ni-enrichment of the austenite particles coupled with completion of
precipitation strengthening for peak aging condition involving
enrichment of the 3 nm Cu precipitates and cementite conversion to
3 nm M.sub.2C carbides. This is achieved by a longer-time final
tempering at a lower temperature characterized by the peak
strengthening condition. Thus, from the toughness-hardness
trajectory for isochronal tempering presented in FIG. 36, the
optimal final stage tempering condition was determined to be about
5 hours at 450.degree. C., which produced a peak hardness of 436
VHN. The first step was optimized by varying the tempering time
from 5 to 90 minutes over a temperature range of 500.degree. C. to
575.degree. C. in intervals of 25.degree. C.
[0159] FIG. 42 shows the variety of two-step heat treatments
investigated to maximize the toughness-strength combination in
comparison with an HSLA100 alloy and is superimposed on the
isochronal tempering plot. The labels in the plot represent the
tempering time in minutes corresponding to the first step and the
bold black arrow points to the condition for maximum strengthening,
which is the final step in the tempering sequence. The short time,
high temperature nucleation treatments were conducted in a molten
salt-bath to reduce heating time followed by water quench to reduce
cooling time. The initial solution treatment was conducted in argon
atmosphere and isothermal aging was conducted under vacuum as
described earlier.
[0160] The optimal combination of toughness and strength is
determined from FIG. 42 to be about a 550.degree. C. 30 minutes
followed by 450.degree. C. 5 hours heat treatment. The apparent
achievement of optimal austenite stability by multi-step tempering
results in significant increase of impact toughness to 130 ft-lbs
at a hardness level of 415 VHN. Comparing with the baseline
toughness-strength combination from isochronal tempering data, a
transformation toughening increment of 50% from 87 ft-lbs for the
10 hour isothermal treatment and 70% from 77 ft-lbs for the 5 hour
isothermal treatment is observed at the same strength level. So an
average of 60% toughness increment due to dispersed phase
transformation toughening can be attributed to multi-step tempering
when compared to standard isothermal tempering at the same strength
level.
[0161] The competition of several substructures begins with the
first higher temperature nucleation treatment. Within the carbide
subsystem, cementite has an initial advantage of precipitation
because it involves only rapid interstitial carbon diffusion. As
aging time increases, the more stable but kinetically slower
M.sub.2C carbides attract carbon from cementite as they coherently
precipitate at heterogeneous sites provided by the high dislocation
density of the martensitic matrix. In parallel, the copper atoms
also partition out of solution and nucleate on the dislocation
substructure. This promotes not only dissolution of cementite but
also heterogeneous nucleation of austenite particles on the carbide
and copper strengthening precipitates. The precipitation phenomenon
is halted after the first step nucleation treatment by water
quenching. At this point, the microstructure consists of embryonic
BCC copper and M.sub.2C precipitates acting as nucleation sites for
intralath austenite with some undissolved cementite. The second
heat treatment step continues the precipitation of M.sub.2C at the
expense of cementite and enriches the fine austenite in Ni while
continuing the precipitation of Cu. The lower temperature of this
second tempering step is likely to produce additional nucleation of
the strengthening precipitates as more dislocation sites are
activated by the higher driving force. The embrittling cementite
dispersion is eventually consumed by the very fine dispersion of
M.sub.2C.
[0162] SEM analysis of the fracture surfaces for the multi-step
treatment specimens indicate transition from quasi-cleavage to
ductile mode of failure as the time of initial tempering is
increased, attributed to transformation toughening increment as
described. FIG. 43 presents a representative micrograph of the
fracture surface for the optimal toughness-strength combination for
tempering treatment of 550.degree. C. 30 min+450.degree. C. 5 hrs.
FIG. 44 shows a higher magnification micrograph of a primary void
in the same sample. The relatively higher degree of primary void
growth is consistent with delayed microvoid instability, as
expected for transformation toughening.
[0163] Mechanical Properties [0164] Evaluation of Tensile
Properties
[0165] An evaluation of the tensile properties was conducted to
determine the actual yield strength of the alloy under the
optimized tempering conditions and to provide a basis for
comparison of the hardness--strength correlation for this class of
steels. Room temperature tensile properties were assessed for the
chosen heat treatment conditions based on the results of the
toughness--hardness data from both isochronal and multi-step
tempering response. The tempering conditions were chosen to cover
the full width of the toughness--strength combination plot (FIG.
42). The same processing route of solution treatment at 900.degree.
C. for 1 hour followed by water and liquid nitrogen quench and
isothermal aging (short time aging was done using molten salt bath)
was followed for the tensile samples, prior to final machining into
dimensions described previously. Duplicate samples for each heat
treatment condition were tested to determine the scatter in the
data. Table 5 summarizes the results of the tensile testing for the
solution treated and aged samples for each heat treatment condition
and provides hardness values for comparison. FIG. 46 presents the
true stress vs. true plastic strain curves for all the samples
tested. The curves are represented as solid lines until the point
of tensile instability (necking) or uniform elongation and by
dotted lines thereafter. The tensile data presented in FIG. 45 and
Table 5 confirms the design of a 160 ksi yield strength steel. The
multi-step tempering treatments helped to achieve the 160 ksi yield
strength goal. TABLE-US-00005 TABLE 5 Room temperature tensile
properties of alloy 0.2% Off-set Yield Ultimate Tensile Uniform
Reduction Strength Strength Elongation in Area Hardness Tempering
Condition ksi (MPa) ksi (MPa) YS/UTS % % VHN 575.degree. C. 5 hours
142.12 (980) 146.45 (2019) 0.97 4.98 73.94 355.30 550.degree. C. 30
minutes + 156.35 (1078) 167.56 (1155) 0.93 5.89 64.60 414.70
450.degree. C. 5 hours 500.degree. C. 30 minutes + 160.97 (1110)
180.16 (1242) 0.89 5.70 57.09 436.57 450.degree. C. 5 hours
[0166] From data on the reduction in area at fracture and uniform
elongation in Table 3, all the heat treatment conditions show
reasonably high values of ductility. The ratio of YS/UTS (strength
ratio) is a general measure of work hardening behavior. The low
values of strength ratio for the "transformation toughening
optimized" multi-step treatments compared to that for the
single-step treatment condition suggests that the work hardening of
the steel is appreciably improved by the optimal tempering
treatments. The load-displacement curves for all the conditions
showed smooth yielding without any distinguishable upper and lower
yield points. Analysis revealed that the plastic stress strain
behavior could be described by the Hollomon power law equation
(Equation 12). The fitting parameters are summarized in Table 6.
.sigma..sub.pl=K.epsilon..sub.pl.sup.n (12) n is the
strain-hardening exponent and K is the strength coefficient in
ksi.
[0167] The yield strength and hardness data from Table 3 is
superimposed on the hardness--yield strength correlation plot in
FIG. 46 (see FIG. 6 for comparison). The black heavy points
represent the data from the current tensile properties study of
alloy. TABLE-US-00006 TABLE 6 Fitting parameters for Hollomon power
law equation from tensile data of alloy (FIG. 45) Strain Strength
Hardening Coefficient, Yield Stress, ksi Exponent ksi Tempering
Condition .sigma..sub.0 n K 575.degree. C. 5 hours # 1 138.59 0.027
162.8 # 2 145.64 0.03 172.6 550.degree. C. 30 minutes + # 1 155.81
0.04 199.1 450.degree. C. 5 hours # 2 156.9 0.038 198.7 500.degree.
C. 30 minutes + # 1 157.5 0.042 216.6 450.degree. C. 5 hours # 2
164.44 0.048 219.8
[0168] Toughness--Temperature Dependence
[0169] To characterize the effect of service temperature on
toughness, Charpy V-notch impact tests were performed over
temperatures ranging from -84.degree. C. to 100.degree. C. for the
tempering condition that optimized the austenite for
room-temperature dispersed phase transformation toughening. Thus,
from FIG. 42 the tempering condition displaying the best
toughness-strength combination, 550.degree. C. 30 min+450.degree.
C. 5 hrs was chosen. The alloys were solution treated at
900.degree. C. for 1 hour, water quenched, liquid nitrogen cooled
and then multi-step tempered. The samples were thermally
equilibrated at the test temperature for 20 minutes prior to
testing.
[0170] FIG. 47 shows the Charpy impact energy of the prototype as a
function of test temperature. The corresponding impact energy
values for 5 hour and 10 hour tempering treatments at room
temperature are superimposed on the plot. The plot shows that there
is a 30 ft-lbs toughness increment at 25.degree. C. compared to the
baseline ductile fracture toughness at lower and higher test
temperatures. Additional toughening occurs in the alloy because of
the delay of microvoid shear localization during ductile fracture
by the optimum stability austenite dispersion. At higher and lower
test temperatures austenite becomes less stable than required for
transformation toughening to occur although the fracture still
occurs in a purely ductile mode, as confirmed by fractography.
[0171] SEM micrographs of the fracture surfaces presented in FIGS.
48 to 52 at each of the test temperatures establish the mode of
fracture. FIG. 48 shows that the fracture surface for the alloy
tested at -84.degree. C. is representative of quasicleavage
fracture characterized by the array of flat facets with dimples and
tear ridges around the periphery of the facets. This indicates a
brittle mode of failure. However, as the test temperature is
increased to -40.degree. C., the fracture surface primarily
consists of microvoids. Although most of the fracture surface is
characteristic of ductile mode of fracture, closer investigation of
FIG. 49 shows that there are a few tear ridges with facets,
indicating a slightly mixed fracture mode. FIGS. 50, 51 and 52 are
representative micrographs from fracture surfaces of alloys tested
at -20.degree. C., 0.degree. C. and 100.degree. C. respectively
showing purely ductile mode of fracture characterized by primary
voids and microvoids without any evidence of flat facets. The
micrographs for the fracture surface of the prototype tested at
room temperature are presented in FIGS. 43 and 44, which contain
mostly primary voids with very few microvoids. The delay of
microvoid shear localization caused by the dispersed phase,
transformation toughened, optimal stability austenite at the
crack-tip stress state leads to more extensive growth of the
primary voids before they coalesce by microvoiding. This finding
further supports transformation toughening by multi-step tempering
to precipitate an optimal stability dispersion of austenite.
Toughness enhancement is increased by a larger volume change. FIG.
47 indicates that the toughness enhancement in the alloy is
30%.
[0172] Microstructural Validation
[0173] Optimization of the processing conditions of the alloy for
dispersed phase transformation toughening in combination with a
fine dispersion of strengthening precipitates has been supported by
property evaluation in the previous sections. Microanalytical
characterization of the austenite dispersion and the strengthening
precipitates and their interaction with the other substructures in
the prototype was performed. [0174] Three-Dimensional Atom Probe
(3DAP) Microscopy
[0175] 3DAP microscopy was chosen to the be the preferred method of
characterization over X-Ray diffraction, Magnetometry and
Transmission Electron Microscopy for identifying the nanometer
scale intra-lath austenite and the optimal 3 nm particle size
strengthening precipitates in the transformation toughened alloy.
This characterization tool was used as a means of evaluating the
matrix composition as well as precipitate compositions, sizes,
morphologies and their average number density.
[0176] The choice of samples for analysis was based on the
condition of tempering treatment for the highest obtainable number
density of the precipitates, determined from the assessed
mechanical properties (FIG. 42). Thus, the tempering condition
corresponding to the highest observed strength (Table 5) namely,
500.degree. C. 30 min+450.degree. C. 5 hrs was chosen. The
450.degree. C. 1 hour tempering condition was also chosen as a
reference for comparison with 3DAP data on similar Cu-strengthened
steels as well as with the other tempering condition. For
simplification, the 450.degree. C. 1 hour tempering treatment
specimen is referred to as the "single-step temper" and the
500.degree. C. 30 min followed by 450.degree. C. 5 hrs tempering
treatment specimen is referred to as the "multi-step temper". The
data for both the tempering conditions is presented simultaneously
for easier comparison.
[0177] The analyzed tips were isothermally aged according to their
respective schedules, following solution treatment at 900.degree.
C. for 1 hour, water quench and liquid nitrogen quench. The overall
composition of the reconstructed volume from atom probe analysis
was obtained and compared with the actual composition of the
prototype as shown in Table 7. It is seen that the actual
compositions compare well with that for the elements detected. The
error for the concentrations is given by 2 .sigma..sub.c, where
.sigma..sub.c= {square root over (c(1-c)/N,)} with c being the
measured composition and N being the total number of atoms
detected. Thus, the statistical error associated with composition
analysis decreases as the total number of atoms detected increases.
TABLE-US-00007 TABLE 7 Comparison between the actual overall
composition of alloy and the overall compositions determined by
3DAP analysis Overall Composition from 3DAP 500.degree. C. 30 min +
Actual Overall Composition 450.degree. C. 1 hr 450.degree. C. 5 hrs
Element wt % at % at % at % Fe 87.2 90 89.90 .+-. 0.08 88.58 .+-.
0.18 C 0.04 0.192 0.11 .+-. 0.24 0.12 .+-. 0.53 Cu 3.64 3.30 2.37
.+-. 0.23 1.13 .+-. 0.53 Ni 6.61 6.49 5.34 .+-. 0.23 7.01 .+-. 0.52
Cr 1.78 1.97 1.86 .+-. 0.23 2.1 .+-. 0.53 Mo 0.58 0.35 0.31 .+-.
0.24 0.89 .+-. 0.53 V 0.11 0.124 0.11 .+-. 0.24 0.16 .+-. 0.53
[0178] Atom probe analysis of the single-step temper was conducted
at 50K while that for the multi-step temper at 70K with a pulse
fraction of 20% at a pulse frequency of 1.5 kHz from 7 kV to 10 kV
steady state DC voltage. The complete analysis for the single-step
temper contained a total of 751,608 atoms in a reconstruction
volume of dimensions 13 nm.times.13 nm.times.84nm. The multi-step
temper analysis collected 254,917 atoms in a reconstruction volume
dimension of 17 nm.times.16 nm.times.28 nm. FIGS. 53 and 54 show
partial 3D reconstruction of all the atoms detected after being
field evaporated from the specimen with their positions and
elemental identities for single-step and multi-step tempering
conditions respectively. Iron is not shown in any reconstruction in
this section for purpose of clarity, enabling larger
microstructural features like precipitates to be seen
distinctly.
[0179] The regions of high copper concentration are clearly
noticeable in both FIGS. 53 and 54 confirming the presence of a
nanometer sized copper particle distribution in the microstructure.
These copper--rich precipitates can be represented by an
isoconcentration surface at 10 at % copper level overlaid with the
atomic positions of copper atoms as shown in FIGS. 55 and 56. The
isoconcentration surfaces clearly outline the Cu-rich precipitates.
The size of the copper precipitates for the single-step temper is
relatively smaller than that for the multi-step temper, while the
number density of precipitates for the former is much higher.
[0180] The shape of the copper precipitates appears to be
elliptical and stretched in the direction of analysis for both the
tempering conditions. The distortion is an instrument artifact due
to a magnification effect caused by the difference in field
evaporation of copper precipitates compared to the matrix. The
precipitates are believed to be spherical in shape.
[0181] Having defined the copper precipitates by the
isoconcentration surface, the size, number densities and
compositions of these copper precipitates can be determined with
the help of the 3DAP analysis software, ADAM. Cross-sectional views
from an analyzed volume of the reconstruction were used to measure
the size of the precipitates. For the single-step temper, the
average diameter of the copper precipitates contained completely
within the analysis volume was found to be 2.67.+-.0.57 nm while
that for the multi-step temper is 3.79.+-.0.13 nm. From the
hardness data, it is apparent that the multi-step temper
corresponds to the peak aging condition. However, considering the
statistical error of the measurement and a distribution of particle
sizes in the material, the optimal particle size of BCC
Cu-precipitates for maximum particle size lies within about 2.5-4
nm.
[0182] The number density of strengthening Cu precipitates is
higher for the single-step temper than the multi-step temper. The
number density of the copper precipitates in the analyzed volume
was estimated by Equation 13. N V = N p .times. .times. .zeta. n
.times. .times. .OMEGA. ( 13 ) ##EQU8##
[0183] N.sub.p and n are the number of particles and the total
number of atoms detected in the volume, .OMEGA. is the average
atomic volume and .zeta. is the detection efficiency of a single
ion detector, equal to 0.6 in this case. The number density of
copper precipitates for the single-step temper was calculated to be
5.42.times.10.sup.18 precipitates/cm.sup.3 while that for
multi-step temper was calculated to be 1.2.times.10.sup.18
precipitates/cm.sup.3. The high number density measured for the
single-step temper (4.5 times that for multi-step temper) is
consistent with the high Cu content of the alloy. Evidence for
cementite dissolution in the toughness-hardness plots of FIG. 42
support the presence of M.sub.2C carbides contributing to the
strength of the multi-step tempered material.
[0184] The average matrix and precipitate compositions can be
determined from the analyzed volume by calculating the fraction of
atoms of each element within the phase. To analyze the composition
of the inner core of the precipitates, a higher threshold level of
15 at % was set to isolate them. Tables 8 and 9 give the
composition of the Cu-precipitates and the matrix respectively with
2 .sigma. a error bar limits for both the single-step and
multi-step conditions. Table 9 also compares alloy matrix
composition with the homogeneous phase composition of the BCC
matrix predicted for austenite stability. TABLE-US-00008 TABLE 8
Average copper precipitate compositions determined by 3DAP analysis
for selected heat treatment compositions. ND means not detected BCC
Cu Precipitate Composition from 3DAP analysis 450.degree. C. 1 hr
500.degree. C. 30 min + 450.degree. C. 5 hrs Element at % at % Fe
30.25 .+-. 3.53 43.79 .+-. 6.52 Cu 63.50 .+-. 2.55 46.69 .+-. 6.35
Ni 5.40 .+-. 4.11 8.76 .+-. 8.31 C ND ND Cr 0.40 .+-. 4.21 0.57
.+-. 8.67 Mo 0.13 .+-. 4.22 ND V ND 0.19 .+-. 8.69
[0185] TABLE-US-00009 TABLE 9 Average matrix compositions
determined by 3DAP analysis for selected heat treatment
compositions compared with equilibrium prediction. ND means not
detected BCC Matrix Composition Equilibrium from 3DAP analysis
Prediction 450.degree. C. 1 hr 500.degree. C. 30 min + 450.degree.
C. 5 hrs 490.degree. C. Element at % at % at % Fe 91.22 .+-. 0.49
92.01 .+-. 0.22 94.1 Cu 0.73 .+-. 0.66 0.22 .+-. 0.77 0.12 Ni 5.32
.+-. 1.62 6.33 .+-. 0.74 3.78 C 0.014 .+-. 1.67 0.041 .+-. 0.77
0.000044 Cr 2.18 .+-. 1.65 0.88 .+-. 0.76 1.88 Mo 0.44 .+-. 1.66
0.39 .+-. 0.77 0.10 V 0.09 .+-. 1.67 0.12 .+-. 0.77 0.02
[0186] The results of the 3DAP analysis indicate that the matrix
composition for both heat treatment conditions compare reasonably
well with the predicted equilibrium calculations. The matrix Cu
composition is near the predicted equilibrium composition at the
earliest evolution stage, indicating a high degree of Cu
precipitation and it remains at the equilibrium condition for the
multi-step temper composition analyzed. The relatively higher Ni
level observed for both conditions may be associated with the
microsegregation compositional banding described earlier. The
difference between the homogeneous equilibrium matrix Ni prediction
and the 3DAP microanalysis results is consistent with the level of
banding microsegregation observed with respect to Ni.
[0187] The average matrix and precipitate compositions and the
concentration of the various solute atoms near the
matrix/precipitate interface can be investigated by a proximity
histogram, or "proxigram", available in ADAM. The concentration
values were determined by averaging the concentration in 0.2 nm
peripheral shells around all the precipitates with respect to the
10 at % copper isoconcentration surface, within and outside the
precipitates. The negative values in abscissa represent the matrix
composition while the positive values are indicative of the
precipitate compositions. However, the zero point is not
necessarily a correct estimate of the precipitate/matrix interface
and serves as an approximate reference point. The proxigrams
obtained from analysis of copper precipitates in single-step temper
and multi-step temper samples are presented in FIGS. 57 and 58
respectively. The proxigrams indicate that for both cases of
tempering condition, Ni shows considerable partitioning to the
precipitate/matrix interface while that for other solute atoms is
within the error limit of estimation. The level of Ni enrichment at
the interface is about 50% higher than the matrix Ni content for
the single-step temper observed in FIG. 57.
[0188] Referring to FIG. 58, that the level of Ni located near the
interface was more than 50% with respect to the matrix Ni content
for the multi-step temper condition. This led to further
investigation of the precipitate/matrix interface region by varying
Ni concentration threshold levels in the 3D reconstruction for the
multi-step temper. Setting a 10 at % level for Ni, the
isoconcentration surface of a Ni-rich precipitate at the interface
of the Cu-rich precipitates could be identified. FIG. 59 shows the
isoconcentration surface outlining the Ni-rich precipitate defined
at 10 at % Ni, overlaid with atomic positions of Cu and Ni from
three different orientations. Composition analysis for the Ni-rich
precipitate and its comparison with equilibrium prediction of
austenite composition is shown in Table 10. Ni concentration of
19.5 at % in the precipitate demonstrates that the precipitate is
the desired austenite of optimum stability for transformation
toughening. Lower than equilibrium concentration of the Ni in the
austenite estimated as 30 at % may be attributed to the local
magnification effects previously mentioned. This is further
supported by the higher (twice) Cu level in austenite than
equilibrium prediction due to the possibility of having copper
atoms from the adjacent copper precipitates projected into the
austenite precipitate because of the solute overlap effect. Since
only a single austenite particle was observed, the statistical
error associated with the composition estimation is likely
significant. To confirm the Ni content of austenite, further
investigation was done by a one-dimensional composition profile
plotted along the atom-probe analysis direction in FIG. 60. This
confirmed that the Ni content of austenite is 30 at % and is
consistent with equilibrium values. TABLE-US-00010 TABLE 10 Average
austenite composition determined by 3DAP analysis for selected heat
treatment compositions compared with equilibrium prediction. ND
means not detected Austenite Composition Equilibrium from 3DAP
analysis Prediction 500.degree. C. 30 min + 450.degree. C. 5 hrs
490.degree. C. Element at % at % Fe 65.9 .+-. 5.6 61.5 Cu 13.9 .+-.
8.9 6.97 Ni 19.3 .+-. 8.6 29.8 C ND 0.00068 Cr 0.93 .+-. 9.6 1.47
Mo ND 0.03 V ND 0.00084
[0189] The size and location of the austenite precipitate, measured
as 5 nm from FIG. 60, confirms that it is intralath austenite
nucleated on two adjacent Cu precipitates. This result provides
direct visual evidence of the heterogeneous nucleation of intralath
austenite on a fine dispersion of strengthening precipitates; Cu
precipitates in this case. This finding also strengthens the
transformation toughening conclusion of an optimal stability
austenite dispersion is effected by employing a multi-step
tempering treatment to nucleate the austenite in the first
tempering step followed by a Ni-enrichment final tempering
step.
[0190] No M.sub.2C carbide precipitate was identified in the
atom-probe reconstructions. Because of the low equilibrium phase
fraction of M.sub.2C calculated for the optimal tempering
treatment, the precipitates might have been excluded from the
analysis volume of the atom probe. Also, detection of carbide
particles is difficult because of differences in the field
evaporation rates between the carbide and the surrounding matrix
that cause the carbide to stick out in relief leading to
tip-fracture. Such a situation was encountered during the
multi-step temper atom-probe run, when a high level of carbon and
molybdenum was observed in the in-situ composition profile during
data collection and the tip fractured soon thereafter. No data
could thus be obtained for 3D reconstruction and characterization
of M.sub.2C carbide in the alloy.
[0191] Impact toughness of 130 ft-lb was achieved at 160 ksi yield
strength for a multi-step tempering condition of the alloy, which
is a significant improvement of properties over other conventional
alloys. FIG. 61 graphically represents the toughness-strength
combination of the alloy for three different tempering conditions
in comparison to other commercial and experimental alloys.
Summary
[0192] To simulate a continuous casting process, a 34 lb (15.4 kg)
Vacuum Induction Melt (VIM) heat of the alloy was slab cast as
1.75'' (4.45 cm) plate, homogenized for 8 hours at 2200.degree. F.
(1204.degree. C.), hot-rolled to 0.45'' (1.14 cm) and then annealed
at 900.degree. F. (482.degree. C.) for 10 hours. Consistent with
microsegregation/homogenization simulations, compositional banding
in the plate was limited to an amplitude of 6-7.5 wt % Ni, 3.5-5 wt
% Cu, 1.6-2 wt % Cr, and 0.2-0.5wt % Mo. Examination of the oxide
scale showed no evidence of hot shortness in the alloy during hot
working. The evaluation of the alloy for different tempering
conditions was conducted under an initial martensitic condition
obtained by austenizing solution treatment at 900.degree. C. for 1
hour followed by a water-quench and a liquid nitrogen cool. Since
this is an alloy for low-cost air-hardenable plate steel,
isothermal transformation kinetics measurements were also
conducted, demonstrating achievement of 50% bainite in 4 minutes at
360.degree. C. Hardness and tensile tests confirmed predicted
precipitation strengthening behavior in quench and tempered
material. Isochronal tempering studies at 1 hour confirmed peak
strengthening at 420.degree. C. with gradual overaging. Multi-step
tempering was employed to optimize the austenite dispersion and a
significant enhancement in toughness was observed with minimal loss
in strength for a 550.degree. C. 30 min+450.degree. C. 5 hrs
tempering condition. An optimal austenite stability was indicated
by a significant increase of impact toughness to 130 ft-lb at a
strength level of 160 ksi. Comparison with the baseline
toughness-strength combination determined by isochronal tempering
studies indicates a significant transformation toughening increment
of 60% in Charpy energy. Tensile tests were conducted on the
preferred tempering conditions to confirm the predicted strength
levels. Charpy impact tests and fractography demonstrate ductile
fracture with C.sub.v>80 ft-lbs down to -40.degree. C., with a
substantial toughness peak at 25.degree. C. Cu particle number
densities and the heterogeneous nucleation of optimal stability
high Ni 5 nm austenite on nanometer-scale copper precipitates in
the multi-step tempered samples were confirmed using three-
dimensional atom probe microscopy. The copper precipitate size was
verified for peak strengthening at about 2-3 nm, and a precipitate
composition of 50-60% copper for short tempering times was
confirmed. The fine austenite dispersion showed a Ni content near
of about 30%.
[0193] Variations of the composition of the steel alloy as well as
the processing thereof may be undertaken without departing from the
spirit and scope of the invention. Therefore, while there have been
described preferred compositions and methods, the invention is to
be limited only by the following claims and equivalents
thereof.
* * * * *