Dynamic Fracture Toughness of High Strength Cast Steels

Paper 12-054.pdf, Page 1 of 17
AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
Dynamic Fracture Toughness of High Strength Cast Steels
L.N. Bartlett, A. Dash, D.C. Van Aken, V.L. Richards, K.D. Peaslee
Missouri University of Science and Engineering, Rolla, Missouri
Copyright 2012 American Foundry Society
ABSTRACT
The dynamic fracture toughness of Cr and Mo steels with
nickel contents of 0, 1.56, and 5.5 wt.% was evaluated
and compared to a lightweight steel of composition
Fe-30.40%Mn-8.83%Al-1.07%Si-0.90%C-0.53%Mo.
Each steel was heat treated to a Rockwell C-scale
hardness range of 36 to 38. The 4130, 4325, and HY130
steels were quench-hardened and tempered. The
lightweight steel was solution treated, water quenched and
age hardened. Of the alloys tested, the lightweight steel,
the 4325 steel and the Al-killed and Ca-treated HY130
steel had similar dynamic fracture toughness values of
153, 153 and 165 kJ/m2, respectively. The 4130 steel had
a much lower toughness of 94 kJ/m.2 The lightweight
Fe-Mn-Al-C alloy performed better at Rockwell C32,
producing the highest measured dynamic fracture
toughness of 376 kJ/m2. Toughness of the Cr and Mo
steels was strongly dependent on deoxidation practice.
Alloys treated with ferro-titanium showed a reduction in
toughness, which was attributed to TiN particles and in
one case eutectic Type II sulfides. Addition of misch
metal to an aluminum and ferro-titanium treated HY130
steel eliminated the Type II sulfides and increased the
dynamic fracture toughness from 58 to 88 kJ/m2. HY130
obtained the highest toughness (165 kJ/m2) when
aluminum deoxidation was followed by calcium
treatment.
INTRODUCTION
The fracture toughness of high strength steel can be
characterized by either a stress intensity or a J-integral
approach to fracture mechanics. For materials with low
toughness where cleavage fracture dominates, there is
little plastic deformation around the crack tip and linear
elastic fracture mechanics (LEFM) dominates.1
Toughness is then evaluated as a critical stress intensity
factor, KIc. For sufficiently ductile materials, such as
austenitic Fe-Mn-Al-C steels, failure is governed by the
flow properties around the crack tip and the LEFM
approach is no longer valid.1 Therefore, for materials
exhibiting crack tip blunting, an elastic-plastic fracture
mechanics (EPFM) approach is required and the fracture
behavior is described by the path independent J integral,
which is equivalent to the energy release rate in elasticplastic materials.1 The fracture toughness of ductile
materials, JIc, is defined as the critical value of J near the
onset of stable crack growth.2
In both of the above determinations of fracture toughness,
the material is assumed to be under quasistatic loading
conditions of less than 2 MPa√m/s.3 A material’s
resistance to fracture is often dependent on the loading
rate; therefore, the static JIc may not be representative of
material behavior at high loading rates. Dynamic fracture
toughness (DFT) under high loading rates is often
difficult to obtain because crack extension during impact
loading is difficult to measure. Instrumented Charpy
impact tests provide a reproducible way of measuring the
time dependency of force and crack displacement at
elevated loading rates and therefore, provide a means of
measuring DFT. Brittle fracture or Type I failure for
linear elastic materials is characterized by crack initiation
at the maximum load, which is followed by unstable crack
propagation to failure. Failure of elastic-plastic materials
is characterized by Type II, Type III, or Type IV
behavior. Type II failure occurs when there is enough
plasticity around the crack tip to allow for a small amount
of stable crack extension before fracture at the peak load.
For structural applications subject to shock loading, Type
III or Type IV failure of the steel is desired. In Type III
failure, there is significant stable crack extension after the
peak load followed by unstable fracture. Type IV fracture
is characterized by stable crack growth and the material
fails by ductile tearing only, which is the desirable
behavior for military armor. Schindler4 has proposed a
method of determining toughness from instrumented
Charpy impact tests that is based on the crack tip opening
displacement (CTOD) and crack tip opening angle
(CTOA) models of crack nucleation and growth. This
results in an algebraic expression for the dynamic J-R
curve from which J1d can be evaluated in an analogous
way to the determination of JIc.2 This method is a single
specimen approach with experimental inputs of the peak
load (Pmax), the energy up to peak load (Emax) and the total
facture energy (Etot), which are easily determined from the
instrumented Charpy results. A complete description of
the test procedure and methods for the analysis of the data
may be obtained from a reading of Schindler4, ASTM E
18203 and ASTM E 813.2
FRACTURE OF HIGH STRENGTH STEELS
The ability of a cast high strength steel to resist fracture is
a function of many different metallurgical factors,
including the inherent matrix toughness, the segregation
of impurities that have limited solubility and the
composition, morphology, and distribution of second
phase particles. For steels of similar strength and
microstructure, fracture toughness at elevated loading
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rates is a strong function of steel cleanliness. Deoxidation
practice can affect the toughness because it affects the
inclusion type, size, morphology and distribution; and,
therefore, can affect the fracture behavior. The use of
strong deoxidizers such as Al and Ti can promote the
formation of eutectic Type II sulfides, which form in
interdendritic regions and reduce both tensile ductility and
notch toughness.5,6 Titanium is often added after Al
deoxidation to reduce nitrogen related gas porosity, but
this practice can promote Type II sulfides. Type I and
Type III sulfides are less detrimental to toughness because
they form at higher temperatures and are less likely to
form eutectic cells along dendritic boundaries during
solidification. Type I sulfides are globular and have the
least effect on toughness while Type III sulfides are often
faceted and these sharp crystallographic features can lead
to void nucleation at small strains. Additions of misch
metal have been shown to promote globular Type I
sulfide formation near the liquidus as well as convert
Type II and Type III sulfides, which have formed by
deoxidation with Al or Ti into Type I sulfides.6 Additions
of Ti and Al are chemically active in the melt and can also
form nitrides during solidification and subsequent
cooling, which can also reduce toughness. Adding excess
Al in amounts more than necessary for deoxidation has
been shown to produce grain boundary precipitation of
AlN, which can lead to “rock candy” fracture.6
Quench-hardened and tempered (Q and T) Cr and Mo
steels are used extensively for structural applications.
Some of the more important variables that influence the
capability of Q and T martensite to suppress brittle
fracture during impact loading are the fineness of the
microstructure and the ability of screw dislocations to
cross slip at high strain rates.5, 7 There is a rich
metallurgical history associated with the addition of Ni to
improve notch toughness. It has been shown that
additions of Ni promotes cross slip of screw dislocations,
produces solid solution softening at low temperature and
lowers the ductile to brittle transition temperature in high
strength steels.5 HY130 steels are high strength Cr and
Mo steels that have been modified with additions of Ni
that typically range between 2 and 5.3 wt.% Ni. 5 All
compositions in the following text are in weight percent.
Yield strength values for Q and T HY130 steels generally
range from 897 to 980 MPa and tensile strengths greater
than 1047 MPa with elongations as high as 19% have
been reported in wrought HY130.5, 8 Dynamic fracture
toughness of Ni modified Cr and Mo steel has been
shown to increase with loading rate and toughness values
greater than 250 kJ/m2 have been reported in wrought
HY80 steel.9
Austenitic steels in the Fe-Mn-Al-C system have gained
much interest as a lightweight alternative to traditional
high strength, cast, and Q and T steels. Depending on
composition, the addition of Al produces a 12 to 18%
reduction in density below that of traditional steel without
a sacrifice in mechanical properties.10-13 Typical alloy
compositions contain from 20-30% Mn, 5-11% Al, and
0.3-1.0% C and are age-hardenable by the coherent
precipitation of nano-sized κ-carbide, (Fe,Mn)3AlC,
which precipitates homogeneously in the austenitic
matrix.11, 14-17
The mechanical properties of Fe-Mn-Al-C alloys depend
on the composition and the degree of age hardening. In a
recent study, Van Aken et al. reported a yield strength of
873 MPa, an ultimate tensile strength of 953 MPa and an
elongation to failure greater than 20% for a cast
Fe-30.40%Mn-8.83%Al-1.07%Si-0.90%C-0.53%Mo
alloy aged to a Rockwell C-scale (HRC) hardness of 38.18
An austenitic matrix and high work hardening rates,
which are characteristic of Fe-Mn-Al-C steels, contribute
to high energy absorption rates under impact loading.
However, the toughness of Fe-Mn-Al-C alloys is a strong
function of P content and P contents greater than 0.01%
have been shown to promote brittle cleavage fracture. 18
Impact toughness has also been shown to be dependent on
inclusion type and density. Schulte et al. have shown that
an increasing amount of AlN precipitation reduced the
breaking energy at -40C (-40F) by almost 50% in a
Fe-30.30%Mn-8.76%Al-1.02%Si-0.94%C-0.38%Mo
alloy.19 However, with a clean melt practice, using high
purity charge materials, and keeping P levels below
0.006%, dynamic fracture toughness values greater than
400 kJ/m2 have recently been reported in age hardened
alloys.20
In the study reported here, the Schindler4 method was
used to determine the dynamic fracture toughness (DFT)
of three different Ni-modified Cr and Mo Q and T steels.
The role of deoxidation practice was also examined and
the results are compared with toughness values obtained
from a lightweight Fe-Mn-Al-C alloy aged to an
equivalent hardness.
DESIGN OF EXPERIMENTS
Five different heats of HY130 were prepared to study the
effects of different deoxidation practices on the same
nominal composition. Three HY130 heats were prepared
at Missouri S and T (HY130 heats A, B and C) and two
heats were prepared at two different commercial
foundries (HY130 heats D and E). Deoxidation practice
for each alloy is listed in Table 1. The furnace charge for
HY130 heats A, B and C consisted of high purity
induction iron, De-sulco graphite, ferro-silicon (75%),
ferro-molybdenum (60%), ferro-vanadium, electrolytic
manganese, electrolytic nickel and electrolytic chromium.
HY130 heats A and B were prepared under Ar cover in
200 lb and 100 lb coreless induction furnaces,
respectively, and poured into bonded silica sand Y-block
molds. HY130 heat C was prepared in a 20 lb vacuum
induction furnace and cast into an investment Y-block
mold. HY130 heats D and E were also prepared in
induction furnaces at commercial foundries.
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Table 1. Deoxidation Practices of The Alloys Used in the Current Study
Alloy/Heat
Quantity/Furnace
Atmosphere
Filter
HY130 - A
140 lb melt in coreless
induction furnace
Ar gas cover
No filters used
HY130 - B
100 lb in coreless
induction furnace
Ar gas cover
64 ppi filter used
HY130 - C
20 lb in vacuum
induction furnace
Vacuum of 1 torr
No filters used
1000 lb in coreless
induction furnace
Inert atmosphere by
SPAL (surface
protective argon
liquid)
No filters used
No protective
atmosphere
No filters used
Ar gas cover
No filters used
Ar gas cover
No filters used
Al addition (0.007% final)
Al addition (0.008% final)
followed by Ca-Si
(0.008% Ca final) and
ferro-titanium (0.044% Ti
final)
Ar gas cover
No filters used
No deoxidation required
HY130 - D
HY130 - E
4325
4130
Fe-Mn-Al-C
200 lb in coreless
induction furnace
coreless induction
furnace (size not
reported)
coreless induction
furnace (size not
reported)
200 lb in coreless
induction furnace
The HY130 castings were double normalized for one hour
at 941 C (1726F) and one hour at 891C (1636F),
austenitized at 841C (1546F) for one hour, and then oil
quenched. The castings were tempered at 611C (1132F)
for one hour to produce similar microstructures and
hardness values. The 4325, 4130, and the lightweight
Fe-Mn-Al-C steels were cast in commercial foundries.
The 4325 alloy was deoxidized with Al and cast into
bonded silica sand tensile bar molds. As cast 4325 bars
were normalized for one hour at 870C (1598F),
austenitized at 875C (1607F) for one hour, water
quenched, and tempered for 30 min at 611C 1132F). The
4130 alloy was deoxidized with Al followed by Ca
addition in the form of Ca-Si, treated with ferro-titanium
additions and cast into bonded silica sand molds.
Specimens were rough cut from the casting and
austenitized at 870C 1598F) for one hour, water
quenched, and tempered for one hour at 480C (896F).
The Fe-Mn-Al-C steel heat was prepared using high
purity induction iron, high purity aluminum, ferro-silicon,
ferro-molybdenum, graphite and electrolytic manganese.
The charge was melted under argon cover, triple calcium
Deoxidation practice
0.07% Al added to
pouring stream in ladle
followed by 0.05% Ca
added as Ca-Si wire
Same as HY130 - A
Melt was poured into
mold directly.
Deoxidation same as
HY130 - A
0.08% Al added in
furnace, followed by
0.12% ferro-titanium,
and 0.08% misch metal.
All additions were added
within 20 s in the
furnace.
Pre-alloyed ingot was
remelted without
additional deoxidation by
Al, and ferro-titanium
was added prior to
pouring (0.11% Ti final).
treated and argon stirred prior to casting into bonded
silica sand plate molds, which were coated with a zircon
wash. Because of the high Al content in Fe-Mn-Al-C
steels, no additional deoxidation practice was required.
Specimens from the Fe-Mn-Al-C plates were solution
treated for two hours at 1050C 1922F) water quenched
and age hardened from 13 to 60 hr at 530C (986F).
Tensile test specimens were prepared from the 4325,
HY130 and Fe-Mn-Al-C steels per ASTM E8 and ASTM
E8M specifications. All tensile specimens were tested on
a MTS servo-hydraulic load frame. DFT specimens were
machined from the center of the castings as standard
rectangular Charpy bars, i.e. 10 mm x 10 mm x 55 mm.
Additional specimens from HY130 heat C were hot
isostatically pressed to determine the effects of porosity
on toughness. A sharp 0.25 mm wide notch with a depth
of 2.0 to 2.5 mm was machined into one side of each test
bar at midspan using a diamond saw (Cr and Mo steels) or
by wire EDM (Fe-Mn-Al-C steel). The notched
specimens were fatigue pre-cracked in 3-pt bending
(R = 0.1) to a total initial crack length (a0) of between 3.5
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AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
RESULTS
and 5.5 mm as per ASTM E1820. Finished DFT bars
were fractured at room temperature on a Tinius Olson
Charpy impact machine outfitted with an MPM
instrumented striker. The load versus displacement data
was used to estimate the dynamic fracture toughness (J1d)
using the single specimen technique developed by
Schindler.4 The force displacement data required
smoothing because of “ringing” oscillations in the results
caused by shear wave propagation and reflection within
the specimen during impact. A customized function for
smoothing of the data that takes into account the
wavelength of the oscillations was used as outlined in
Kalthoff and Gregor.21
CHEMICAL ANALYSIS
Chemical analysis for each steel is listed in Table 2. The
oxygen and nitrogen content of the five different HY130
heats is given in Fig. 1. The nitrogen content of the
HY130 steels varied between 40 and 190 ppm. As
expected, the vacuum melted HY130 heat C showed the
least amount of total N content. HY130 heat D proved to
have the lowest O content.
The higher oxygen content in HY130 heat C is attributed
to the presence of moisture in the ceramic investment
mold, which was fired but not pre-heated prior to casting.
The low oxygen content of HY130 heat D was attributed
to the surface protective argon liquid (SPAL) practice
used during melting. The lack of protective atmosphere
during melting of HY130 heat E resulted in the highest
total N and O contents of 190 and 350 ppm.
Chemical analysis of the castings was determined by
using arc spectroscopy for the Cr and Mo steels and by
inductively coupled plasma spectrometry (ICP) after
sample dissolution in perchloric acid for the Fe-Mn-Al-C
steel. Additionally, the oxygen and nitrogen content in
the castings was measured using a LECO O-N analyzer.
Fractography was performed on the ends of broken test
specimens using a Hitachi S570 scanning electron
microscope (SEM). Inclusion analysis was determined
using an ASPEX PICA 1020 SEM with automated feature
analysis. Reported uncertainties were calculated as
sample standard deviation for a sample size greater than
three.
INCLUSION ANALYSIS
The inclusion density by type for steels without Ti
treatment is shown in Fig. 2. All of the inclusion analyses
for the different steels were determined for a particle or
pore major diameter of 1 to 80 µm. Cr and Mo steels
without ferro-titanium additions consisted mainly of
calcium aluminate type inclusions and sulfides containing
Ca and Mn. HY130 heat A, which was prepared using Ar
cover and cast without a filter, had the largest population
of aluminate and sulfide inclusions. Analysis of the
Fe-Mn-Al-C steel showed that inclusions were mainly
complexes of AlN and MnS. In many cases, MnS was
Table 2. Chemical Analysis of Different Steel Castings Given in Weight Percent
Alloy/Heat
C
Si
Mn
P
S
Cr
Mo
Ni
Al
Ti
V
Fe
HY130 / A
.183
.53
.69
<.0012
.0014
.76
.51
4.7
.09
-
.08
bal
HY130 / B
.18
.56
.73
<.0012
.004
.57
.52
5.37
.064
.002
.063
bal
HY130 / C
.18
.70
.54
<.0012
<.007
.60
.49
5.07
.023
-
.074
bal
HY130 / D
.17
.47
.77
<.0012
.001
.71
.50
>6.0
.028
.044
.08
bal
HY130 / E
.21
.23
.41
<.0012
<.007
.663
.67
5.4
.009
.11
.145
bal
4325
.28
.82
.63
.0077
.008
.82
.29
1.56
.007
0.014
0.005
bal
4130
.35
.46
.74
<.0012
<.007
1.02
.24
.008
.008
.044
<.002
bal
Fe-Mn-Al-C
.90
1.07
30.42
.001
.006
-
.53
.008
8.83
-
-
bal
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Fig. 1. The oxygen and nitrogen contents of the
HY130 heats show that the highest amount of both
oxygen and nitrogen were obtained for HY130 heat E
which was melted without protective atmosphere.
Fig. 3. The inclusion density by type for the alloys
with ferro-titanium addition shows that the largest
percentage of the inclusions in the Ti treated HY130
and 4130 alloys were TiN and MnS. The HY130 heat E
steel, which was melted without protective
atmosphere, had more than twice the amount of
inclusions as the other steels.
heat E, which was prepared without protective
atmosphere, had more than twice the amount of TiN and.
MnS inclusions as the 4130 and the SPAL processed
HY130, heat D.
Fig. 2. The inclusion density by type for the alloys
without Ti addition shows that a large percentage of
the inclusions in the HY130 and 4325 alloys were
calcium aluminate and (Mn, Ca)S inclusions.
Inclusions in the Fe-Mn-Al-C steel were mostly
complex AlN and MnS.
found to precipitate around pre-existing AlN inclusions.
The Fe-Mn-Al-C steel also contained the most porosity.
Of the HY130 heats analyzed, heat C was found to
contain the most porosity with a volume fraction of
1.8×10.-4 The inclusion density by type for Cr and Mo
steels with ferro-titanium treatment is shown in Fig. 3.
Most of the inclusions in all three Ti-treated steels
consisted of TiN and MnS with a small percentage of
aluminate inclusions. In many cases, MnS was found to
have precipitated on pre-existing TiN inclusions. HY130
The total average inclusion density was determined for
each steel and the results are given in Fig. 4a. HY130
heat A, which was Al and Ca treated and cast without a
filter, and HY130 heat E, which was Ti-treated and
prepared without a protective atmosphere, had the greatest
inclusion densities at 267 and 247 particles/mm,2
respectively. In fact, both HY130 heats A and E
contained more than twice the amount of inclusions as the
other steels. The 4325 steel had the least amount of
inclusions at 46 particles/mm.2 The HY130 heats D and
C had the largest average inclusion size of greater than
3 μm (Fig. 4b). This is an interesting result since HY130
heats D and C were produced under SPAL and a vacuum,
respectively. The smallest (< 2.5 μm) inclusions were
found in the HY130 heat A and Fe-Mn-Al-C steels. It
should be noted that direct correlation between nitrogen
and oxygen contents (Fig. 1) with inclusion measurements
(Fig. 4) is not possible since inclusions with diameters
less than 1 µm are not included in the later. The total
average volume fraction of inclusions was determined for
the different steels and the results are given in Fig. 4c.
The reported inclusion volume fraction was obtained
directly from the inclusion area fraction coverage. HY130
heat E, which was Ti treated and melted without
protective cover, had the largest volume fraction of
inclusions with a value of 1.9×10.-3 The 4325 was shown
to have the least volume fraction of inclusions at 2.9×10.-4
The volume fraction of inclusions was statistically the
.
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(a)
(b)
(c)
(d)
Fig. 4. (a) The total inclusion density as a function of alloy type and heat number shows that HY130 heat A, which
was prepared under Ar cover and cast without a filter, and HY130 heat E, which was Ti treated and prepared without
protective atmosphere, have more than twice the amount of inclusions as the other steels. (b) HY130 heat D, which
was Ti treated and melted under SPAL, had the largest average inclusion size while HY130 heat A showed the
smallest average inclusion size. (c) HY130 heat E had the largest volume fraction of inclusions and the 4325 showed
least amount. (d) The average inclusion spacing was the highest for the 4325 and the lowest for Al and Ca treated
HY130 heat A, which was cast without a filter.
same for the other alloys regardless of deoxidation
practice or chemistry. Average inclusion spacings for the
different steel castings are shown in Fig. 4d and are
reported as a combination of both non-metallic inclusions
as well as voids and pores. Average inclusion spacing
was determined directly from the automated feature
analysis data, where the closest neighbor distance was
determined from the Cartesian coordinates of each
measured particle.
The 4325 steel had the largest spacing with an average
distance of 63 µm and HY130 heats A and E had the
smallest inclusion spacing with values of 29 and 30 µm,
respectively
METALLOGRAPHY
All of the Cr and Mo steels had similar microstructures of
tempered martensite with an average martensite block
width of 3.5 to 6 μm. Optical micrographs of HY130
heats B and D are shown in Figs.5a and b. Both show
similar microstructures and fineness of the tempered
martensite. Large TiN inclusions, greater than 5 µm in
diameter, were observed in the microstructure of the
HY130 heat D specimen (Fig. 5b). The 4325 and 4130
steel are shown in Figs. 5c and 5d and are similar to the
microstructures of the HY130 steels. The microstructures
of the Fe-Mn-Al-C specimens are shown in Figs. 5e and
5f. The Fe-Mn-Al-C specimen was aged for 13 hr at
530C (986F) and shows a fully austenitic matrix with
carbide and or intermetallic phases that have precipitated
on grain boundaries. The Fe-Mn-Al-C alloy that was
aged for 60 hr at 530C (986F) also shows precipitation on
grain boundaries as well as approximately 1% of ferrite as
shown in Fig. 5f.
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TiN
(a)
(b)
(c)
(d)
Ferrite
(e)
(f)
Fig. 5. Optical micrographs of the (a) HY130 heat B, (b) HY130 heat D, (c) 4325, and (d) 4130 steel specimens show
similar microstructures of tempered lath martensite. (c) A single TiN is shown in the HY130 heat 4 specimen. The
optical micrographs of the Fe-Mn-Al-C alloy (e) aged at 530 C for 13 hr and (f) aged at 530C for 60 hr have been etched
to reveal the dendritic pattern of the austenitic matrix. Small amounts of ferrite (< 1%) were present in the 60 hr aged
specimen in (e). Small angle grain boundaries are highlighted by precipitation of carbides or intermetallic
compounds in the Fe-Mn-Al-C alloys. All specimens were etched with 2% nital followed by etching with Klemms’s
reagent.
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TENSILE PROPERTIES
Tensile properties were evaluated from select heats. The
results are listed in Table 3. Yield and ultimate strengths
of the HY130 specimens were shown to be insensitive to
deoxidation practice and both the yield strength (YS) and
ultimate tensile strength (UTS) were statistically the same
at 970 and 1180 MPa, respectively. Total elongation to
failure (% ef) and percent reduction in area (% RA) was
lowest in HY130 heat E, which was Ti treated and
produced eutectic Type II sulfides. Yield and ultimate
tensile strengths of the 4325 alloy were lower than the
HY130 specimens at 957 and 1094 MPa, respectively.
Elongation to failure was nearly the same between the Al
and Ca deoxidized HY130 heat A castings and the 4325
castings at 11%. However, the HY130 heat A specimen
showed the greatest reduction in area at 43% compared to
the 4325 steel with a reduction in area of 25%.
Tensile properties of the lightweight and austenitic FeMn-Al-C steel are also shown in Table 3 for two different
aged conditions. Increasing the aging time from 26 hr to
60 hr at 530C (986F) increased the hardness from 32 to
38 HRC, increased the yield strength from 728 to 873
MPa and increased the ultimate tensile strength from 873
to 953 MPa. However, the elongation decreased from 28
to 20% with increased aging. At equivalent hardness
values, ie. 36 to 38 HRC, the Fe-Mn-Al-C steel showed
the greatest elongation to failure. The large cast grain
structure of the Fe-Mn-Al-C steel produced an irregularshaped necked region that was difficult to accurately
evaluate for reduction of area measurements and was thus
not reported.
DYNAMIC FRACTURE TOUGHNESS
Select load versus displacement curves from the
instrumented Charpy impact tests are given in Fig. 6 for
the various steels. Force data was smoothed using a sixth
order polynomial fit to eliminate the specimen “ringing”
and to calculate the dynamic fracture toughness. The
HY130 heat B, 4325, and both of the Fe-Mn-Al-C steel
specimens show Type IV behavior with stable crack
extension to failure. The 4130 steel exhibited Type III
behavior with stable crack propagation followed by a drop
in the load versus displacement curve, which is indicative
of unstable crack propagation, subsequent crack arrest and
continued stable crack growth (Fig. 6). The HY130 heat
E showed brittle Type I behavior with unstable crack
propagation and failure of the specimen soon after
reaching the peak load (Fig. 6b). The total energy of the
fracture process for the representative specimens is also
given in Fig. 6 as a function of crack displacement. The
Fe-Mn-Al-C steel specimens that were aged for 13 hr at
530C (986F) to a hardness of 32 HRC had the greatest
total energy at fracture as shown in Fig. 6e. Upon further
aging of the Fe-Mn-Al-C steel to a hardness of 38 HRC,
the absorbed energy of fracture decreased (Fig. 6f) and
the alloy displayed similar fracture energies as the Al and
Ca treated HY130 heat B (Fig. 6a) and the Al killed 4325
steel as shown in Fig.6c.
The dynamic fracture toughness values of the respective
steels are given in Table 4. Of the steels heat treated to
similar hardness values of 36 to 38 HRC, the HY130 heat
B obtained the highest average toughness of 165 kJ/m2.
The 4325 and the Fe-Mn-Al-C steel (aged for 60 hr at
530C [986F)) had similar toughness values of 153 kJ/m2.
The Fe-Mn-Al-C steel achieved a much higher toughness
of 376 kJ/m2 when aged to a hardness of 32 HRC (13 hr
at 530C [986F]).
Of the Al killed and Ca treated, HY130 steels, heat C
proved to have the lowest toughness of only 114 kJ/m.2
Hot isostatic pressing (HIP) of HY130 Heat C specimens
resulted in a 20% reduction in the amount of porosity, a
46% reduction in the size of the remaining voids and
recovery of toughness to an average value of 162 kJ/m.2
The HY130 and 4130 steels, which were Ti-treated,
showed a severe decrease in toughness as compared to
steels deoxidized with Al only or by a combination of Al
and Ca. The HY130 heat E steel, which was Ti treated
and contained eutectic Type II sulfides had the lowest
toughness of only 59 kJ/m2.
FRACTOGRAPHY
The fracture surface of the HY130 heat B specimen is
shown in Fig. 7 a. Failure was mostly ductile in nature
with microvoid initiation and coalescence around globular
calcium aluminate and (Ca, Mn)S inclusions. Additions
of ferro-titanium to the HY130 steels produced large,
> 7 µm, TiN inclusions that are shown fractured on the
surface of the broken specimen from HY130 heat D
(Fig. 7b). In many cases, voids were shown to nucleate
from the fractured TiN inclusions. In one instance, a
quasi-cleavage crack initiated from the fractured TiN
particle as shown in Fig. 7b. The lack of sustained
fracture by cleavage would indicate that the failure was
not hydrogen related
Table 3. Tensile Properties of the Various High Strength Steel Alloys
Alloy/Heat
Heat Treatment
YS, MPa
UTS, MPa
% ef
% RA
2
HY130 - A
Q&T
970 ± 5
1180 ± 5
11.5 ± 2
43 ± 9
2
HY130 – E
Q&T
975 ± 5
1180 ± 5
5.7 ± 2.6
15 ± 1
3
4325
Q&T
957
1094
11.2
25
1
Fe-Mn-Al-C
Aged 26 hrs @ 530° C 728 ± 13
795 ± 11
27.7 ± 6.6
1
Fe-Mn-Al-C
Aged 60 hrs @ 530° C 873 ± 17
953 ± 25
20.1 ± 10.1
1
Elongation measured in a 25 mm gage, 2 Elongation measured in a 30 mm gage, 3 Elongation measured in a 50 mm gage
Hardness,
HRC
36 ± 1
36 ± 1
37 ± 1
32 ± 1
38 ± 1
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(a)
(b)
(c)
(d)
(e)
(f)
Fig. 6. The force vs. displacement curves obtained from the instrumented Charpy tests for the (a) HY130 heat B,(b)
HY130 heat E, (c) 4325, (d) 4130, (e) 13 hr aged Fe-Mn-Al-C steel and (f) 60 hr aged Fe-Mn-Al-C steel show Type IV
behavior for the HY130 heat B, 4325, and both aged Fe-Mn-Al-C steel specimens. The HY130 heat E specimen in (b)
shows brittle Type I behavior and failure soon after peak load. The 4130 specimen in (d) exhibits type III behavior.
The 13 hr aged Fe-Mn-Al-C specimen had the greatest total fracture energy but the lowest peak load.
Paper 12-054.pdf, Page 10 of 17
AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
Table 4. Dynamic Fracture Toughness of the High Strength Steels
Alloy/Heat
Deoxidation
practice
Hardness
(HRC)
Avg. JId
2
(kJ/m )
Inclusion
density,
2
#/mm
HY130 heat A
Al + Ca
37.7 ± .9
122 ± 5.7
267
HY130 heat B
Al + Ca
36.9 ± 1
165 ± 2
115
HY130 heat C
Al + Ca
36.2 ± .2
114 ± 8
HY130 heat C, HIP
Al + Ca
36.1
162
HY130 heat D
Al + Ti + RE
37.7 ± .8
88 ± 9
78
HY130 heat E
Ti
37.4 ± 1
59 ± 4
247
4325
Al only
37 ± 2
153
46
4130
Al + Ca + Ti
38 ± 1
94
121
Al
32 ± 1
376 ± 69
Al
38 ± 2
153 ± 25
76
Fe-Mn-Al-C aged 13 hrs @
530°C
Fe-Mn-Al-C aged 60 hrs @
530°C
77
The HY130 heat D alloy also contained a large amount of
globular MnS inclusions. These globular MnS inclusions
failed by decohesion of the matrix resulting in ductile
tearing (Fig. 7b). In the HY130 heat E specimen, fracture
was also shown to initiate at cracked TiN inclusions.
However, a great deal of the fracture surface showed
failure facilitated by the presence of large, > 50 µm,
eutectic cells of Type II MnS inclusions on prior dendritic
boundaries (Fig. 7c). The fracture surface of one of the
4325 specimens is shown in Fig. 7d and shows ductile
fracture around globular MnS inclusions. Fracture of the
4130 steel specimen revealed the presence of a high
density of cracked TiN particles that were closely spaced
to one another as shown in Fig. 7e. No evidence of
embrittlement from tempering was observed in the 4130
steel despite the unorthodox tempering temperature used.
As in the Ti treated HY130 specimens, fracture was
observed to have initiated in the TiN particles. The
fracture surface of the austenitic Fe-Mn-Al-C steel
specimen, which was aged to a hardness of 38 HRC, is
shown in Fig. 7f. Failure of the Fe-Mn-Al-C specimen
was ductile in nature with microvoid nucleation around
globular MnS inclusions or AlN particles.
DISCUSSION
TENSILE PROPERTIES
Yield and tensile strengths of the Cr and Mo steels were
relatively uniform whereas deoxidation practice had a
significant influence on ductility. Total elongation to
failure decreased by 50% and percent reduction in area
decreased by almost 66% in HY130 heat E when
compared to HY130 heat A. The total inclusion density
between these two steels was largely the same at 267
particles/mm2 for the HY130 heat A and 247
particles/mm2 for the HY130 heat E. However, additions
of ferro-titanium in the HY130 heat E produced large,
greater than 5 µm, TiN inclusions and eutectic Type II
sulfides that contributed to reduced ductility. Strength
and elongation values for the Fe-Mn-Al-C steel were in
good agreement with values reported by Van Aken et al.
for similar compositions of alloys aged to equivalent
hardness levels.18
DYNAMIC FRACTURE TOUGHNESS
Fracture in the high strength steels was dependent on the
nature of second phase particles within the microstructure.
During fracture, void nucleation can occur at the
particle/matrix interface by decohesion or can be created
at brittle particles that fracture during deformation.
Failure then progresses by the growth of voids due to
further plastic strain and then linkage of the voids leading
to final fracture.1 In the presence of a sharp crack or a
notch, crack progression is often dependent on inclusions
and second phase particles that fracture or nucleate voids
within the crack tip process zone.22 In the case of the
Ti-treated materials it is apparent that under dynamic
loading the fracture strength of the TiN is less than the
cohesive strength of the particle and matrix interface. In
contrast, the calcium aluminate particles fail by
decohesion, which indicates that the fracture strength of
the calcium aluminate particle is greater than the cohesive
interface strength.
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AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
(a)
(b)
(c)
(d)
(e)
(f)
Fig. 7. (a) The fracture surface of a specimen from HY130 heat B shows ductile failure with microvoid nucleation and
coalescence of voids around globular calcium aluminate and (Ca, Mn)S inclusions. (b) The surface of the HY130 heat
D specimen shows void nucleation initiated at fractured TiN inclusions. A singular and isolated area of quasicleavage fracture which initiated from the corner of a coarse (>7 µm) TiN particle is highlighted by the arrow in (b).
(c) Fracture in the HY130 heat E specimen was initiated both by fractured TiN and by the presence of Type II MnS
stringers in interdendritic regions. (d) The 4325 specimen shows ductile failure around globular MnS. (e) The
fracture surface of the 4130 alloy shows fracture through a high density of TiN particles which are closely spaced. (f)
Failure in the Fe-Mn-Al-C steel was mostly ductile in nature with microvoid nucleation around MnS and AlN particles.
Paper 12-054.pdf, Page 12 of 17
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The dynamic fracture toughness of the high strength steels
in the current investigation was a strong function of
inclusion content, morphology and spacing. Heats
deoxidized with Al or a combination of Al and Ca
contained mostly globular calcium aluminate and MnS
inclusions that facilitated void rupture at the particle
interface as shown in Figs. 7a and 7d with Type IV
elastic-plastic behavior (Figs. 6a and 6c) and a higher
toughness than heats treated with ferro-titanium.
Toughness of the Al and Ca treated HY130 and 4325
castings varied between 114 and 165kJ/m2, which is much
higher than reported literature values of 92 kJ/m2 for a
wrought 4340 steel.23 It was determined that the
uncharacteristically low toughness of HY130 heat C,
which was vacuum induction melted, was due to the
presence of a high volume fraction, 1.8×10,-4 of porosity.
Hot isostatic pressing was effective at reducing the
volume fraction of porosity from 1.8×10-4 to 1.0×10-5 and
the average pore/void size from 2.7 to 1.4 µm which
resulted in a restoration of fracture toughness to 162
kJ/m2, which was statistically equivalent to the toughness
value of 165 kJ/m2 that was observed from HY130 heat
B. It should be noted that all of the cast steels had
microporosity on the order of 1.0x10.-5 Ferro-titanium
treatment produced large and widespread precipitation of
TiN in HY130 and 4130, which contributed to Type I,
failure in HY130 heat E and Type III failure in the 4130
steel with an accompanying loss in toughness (Figs. 6b
and 6e). In addition to the steels in this study, coarse TiN
precipitation has been shown by several researchers to
lead to brittle fracture and reduced toughness in HSLA
(high strength low alloy) steels as determined by
instrumented Charpy testing.24-26 The deleterious effects
of TiN during fracture was also supported by the
fractography that revealed large void formation from
fractured TiN (Fig. 7b). Fracture toughness was lowest in
the HY130 Heat E steel and this was attributed to a large
concentration of TiN particles as well as Type II MnS
inclusions, which produced unstable crack growth along
the eutectic cells (Fig. 7c). It is well known that without
rare earth additions, the toughness of Cr and Mo steels is
a strong function of S content because of the tendency to
form eutectic Type II MnS inclusions in the vicinity of
grain boundaries.5, 27 Strong deoxidizers such as Ti can
magnify this effect.5, 6 Addition of misch metal to the
ferro-titanium treated HY130 heat D produced globular
MnS inclusions and improved the toughness from 59 to
88 kJ/m.2 However, the presence of TiN appears to
adversely affect void nucleation by particle cracking
rather than interface void nucleation and as a result
produce larger initial voids with commensurate loss in
toughness.
No evidence of hydrogen embrittlement could be
discerned for the Ti-treated steels. In fact, in a study by
Kuslitskii et al.28 on a plain carbon steel with different
types of nonmetallic inclusions, H-embrittlement was
found to be mainly a function of surface area of the
inclusion/metal interface rather than chemistry. Elongated
silicate and alumina inclusions trapped more H and
contributed to a reduction in notched tensile strength.
TiN particles were the least efficient at trapping H and the
TiN containing steels showed the least susceptibility to
H-embrittlement.28
High work hardening rates associated with high Mn-C
austenitic steels have been shown to produce excellent
Charpy V-notch (CVN) toughness. In the solution treated
condition and at room temperature, Hale and Baker
recorded a CVN breaking energy of 206 J for a wrought
Fe-30Mn-8Al-1C alloy.29 Precipitation of κ-carbide
during aging decreases notch toughness and recent studies
of the current Fe-Mn-Al-C steel gave an average value of
115 J for the room temperature breaking energy of a
casting aged in the hardness range of 29 to 32 HRC. 20
The dynamic fracture toughness of the lightweight
Fe-Mn-Al-C steel was reported to exceed that of the Q
and T Cr and Mo steels.20 Fe-Mn-Al-C specimens that
were aged for 13 hr at 530 C (986F) to a hardness of 32
HRC obtained much higher toughness, 376 kJ/m2, as
compared to the Q and T Cr and Mo steels. However,
with an increase in aging to a hardness of 38 HRC, the
Fe-Mn-Al-C steel attained statistically equivalent DFT
values, 153 kJ/m2, as the Al and Ca treated HY130 and
4325 steels. Inclusions present in the Fe-Mn-Al-C steel
were multiphase inclusions composed mainly of AlN and
globular MnS. In most cases, the faceted AlN particles
were coated with globular MnO or MnS that prevented
early void nucleation at the AlN corners or possible
particle fracture.
An attempt was made to model the effects of inclusion
density, size, volume fraction and spacing on the dynamic
fracture toughness of the high strength steels studied here.
The energy of fracture in high strength steels is due in
part by the size of the plastic zone and the type and
distribution of second phase inclusions.22 The plastic
zone size increases with applied stress and is controlled
by the yield strength and fracture toughness. Crack
propagation occurs when the plastic zone size equals the
spacing of the second phase inclusions. Large plastic
zones in front of the crack tip increase the chance of
fracture taking place by direct linkup of voids produced
by fibrous rupture of the matrix. As the yield strength of
the material decreases, the size of the plastic zone
increases and this correlates with an increase in
toughness.22 When the inclusion spacing in the matrix is
large, the interaction between the plastic zones of the
individual particles is limited and this will permit large
strains to occur before void coalescence. Low energy
ductile fracture can occur when there is a high density of
fractured particles with small spacing between them. A
number of researchers have attempted to model the effect
of second phase particles on ductile rupture of both
ferrous and non-ferrous alloys. McClintock30 and
Beremin31 assumed that fracture occurred when the
volume fraction of cavities reached a critical value.
Broek32 studied the role of second phase particles in
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AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
initiating microvoids in a variety of Al alloys and
determined that larger particles with greater inclusion
spacing produced considerable void growth before failure
while smaller and closer spaced particles nucleated
smaller voids, which grew quickly to failure. Broek32
therefore suggested that the fracture strain was inversely
proportional to some function of the volume fraction of
particles or voids and that toughness would be expected to
increase with decreasing particle density. Observations
by Deimel et al. on the relationship of inclusion density to
the fracture toughness of high Cr and Mo pressure vessel
steel showed that fracture toughness was inversely
proportional to the total inclusion density.33 A plot of
inclusion density versus dynamic fracture toughness for
the steels in the current study is given in Fig.8.
regular array of spherical particles with
L0 = D0 (π/6)1/3Vf.-1/3 The fracture toughness was shown
to fit the following model.
KIC ~ Vf -1/6 [2(π/6)1/3YE D0]1/2
Equation 1
Where Y is the tensile yield strength and E is Young’s
modulus. Prior researchers 36, 37have also shown that the
fracture toughness of high strength steel varied linearly
with
An attempt was made to fit the above model
which relates the quasi-static fracture toughness to
inclusion size and distribution to the dynamic fracture
toughness of the steels studied here. The relationship
between inclusion spacing and the volume fraction of
inclusions is given in Fig. 9 and shows an increase in the
spacing between inclusions with decreasing volume
fraction. The slope of the regression line in Fig. 9 gives a
value of -0.37, which is in reasonable agreement with the
value of -0.33 given by Hahn et al.34, 35 For application of
this model to the current study, KId must be converted to
an appropriate JId value by use of the following
equation:1, 33
KId = (EJId)1/2
Equation 2
Substituting Equation 2 into Equation 1 yields the
following relationship between JId and the volume fraction
of inclusions.
JId = Vf -1/3 [2 (π/6)1/3YD0]
Equation 3
Fig. 8. Inclusion density was found to have an inverse
relationship with toughness.
Two distinct trends are observed that suggest that void
nucleation by particle cracking (Ti treated steels) is
inherently worse than void nucleation (Ca and Al treated
steels) at the particle and matrix interface. The fracture
toughness of both the Cr and Mo steels, as well as, the
lightweight Fe-Mn-Al-C steel show a decrease in
toughness with increasing inclusion density (Fig. 8). The
HY130 heat B steel showed a 10 kJ/m2 increase in
toughness over the 4325 steel even though the HY130
steel had more than twice the total inclusion density as the
4325, suggesting a beneficial effect on toughness from the
additional Ni content. However, Ni was shown to have
no noticeable effect on toughness in the Ti treated steels
and the 4130 steel obtained the highest toughness among
the Ti treated steels at 94 kJ/m2. This suggests that in the
Ti treated steels, the favorable effect of Ni at toughening
the matrix is outweighed by the deleterious effect of TiN
particle fracture.
Hahn et al. modeled the effect of inclusion diameter (D0),
spacing (L0), and volume fraction (Vf) on the plane strain
fracture toughness, KIC, of aluminum alloys.34, 35 The
model by Hahn et al. uses an ideal case consisting of a
Fig. 9. Inclusion spacing decreased as the volume
fraction of inclusions increased. The slope of the line
gives a value of -0.38 which is comparable to the
34, 35.
expected value of -0.33.
Figure 10 shows the dependence of dynamic fracture
toughness on the volume fraction of inclusions for the Cr
and Mo steels as well as the Fe-Mn-Al-C alloy. A linear
relationship between Vf -1/3 and DFT is hard to claim for
the steels that were deoxidized with Al or a combination
of both Al and Ca. In fact, DFT was relatively insensitive
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AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
to the volume fraction of inclusions for all of the Al and
Ca treated steels. The steels that were treated with ferrotitanium, however, show that DFT may vary linearly with
Vf,-1/3 but additional testing would be required to verify
the trend (Fig. 10). At the same nominal volume fraction
of inclusions, the Ca and Al treated steels obtained much
higher toughness than the steels treated with Ti. Again,
this may be related to particle cracking in the Ti treated
steel versus interface void nucleation for Al and Ca
treated steel.
Fig. 10. DFT was observed to be a linear function of
-1/3
Vf in the Ti treated steels. At equivalent volume
fractions of inclusions, the Al and Al + Ca treated
steels showed much higher toughness than the steels
which were treated with Ti.
with Ti. This decrease in toughness can again be
attributed to the nucleation of large voids resulting from
TiN particle fracture (Fig. 7b). Conversely, in the steels
which were deoxidized with Al or a combination of Al
and Ca, voids were nucleated by decohesion of the matrix
from the mostly globular calcium aluminate or MnS
inclusions which did not crack during dynamic loading.
Fig. 11. Both the Al and Al + Ca treated castings show
a linear increase in toughness with increasing
inclusion spacing, except for the 4325 casting which
did not fit the expected trend. The castings that were
treated with ferro-titanium also showed a linear trend
but at lower toughness.
From the model by Hahn et al., fracture toughness (KIc) is
a parabolic function of the inclusion spacing as given in
the following relationship.34
KIC ~ [2EYL0]1/2
Equation 4
Combining Equations 2 and 4 gives a linear relationship
between JId and the average inclusion spacing.
JId ~ 2YL0
Equation 5
Figure 11 shows a general trend between increasing
particle spacing (L0) and increasing DFT for all of the Al
and Al plus Ca treated castings and possibly the Ti treated
steels albeit lower.
According to Equation 3, the relationship between D0 and
JId should also be linear. A plot of average inclusion
diameter versus DFT for the high strength steels is shown
in Fig. 12. Toughness was found to increase linearly with
D0 for the Al and Ca treated steels. However, while there
is a general increase in toughness with increasing particle
size in the Ti treated heats, a direct and linear relationship
is difficult to claim. At an equivalent average inclusion
diameter, D0, the Al and Ca deoxidized steels obtained
almost twice the toughness of the steels that were treated
Fig. 12. Although toughness increased linearly with
increasing average inclusion diameter for the Al and
Al + Ca deoxidized castings, a linear relationship is
less obvious in the ferro-titanium treated castings. At
a similar average inclusion size, the steels which were
2
treated with ferro-titanium showed up to a 100 kJ/m
loss in toughness when compared with steels which
were deoxidized with Al or a combination of both Al
and Ca.
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In sufficiently ductile materials, fracture is the
combination of void nucleation (by particle/matrix
decohesion or particle cracking), void growth and
ultimate coalescence to failure. The true fracture strain is
then a function of the strain to nucleate voids, εn, and the
strain to grow them to failure, εg.
εfracture = εn + εg
Equation 6
In general, εn increases if particles are small and round,
but decreases when particles are faceted, brittle and
clustered. The Brown-Embury38 equation relates the
strain associated with void growth to the volume fraction
of inclusions as seen in the following:
[(
)
( ) ]
Equation 7
True fracture strain as determined from tensile testing was
used along with the Brown-Embury equation to determine
the strain to nucleate, εn, for the Al and Ti treated steels.
However, all of the calculated values of εn were negative,
suggesting that the strain to fracture was almost entirely a
function of the strain for void growth for the steels in the
present study.
Deoxidation with Ti produced a large density TiN
particles that cracked during deformation producing large
voids, which decreased the distance between major voids
and reduced the toughness. A similar observation for the
toughness was observed between HY130 heat C HIP and
the as cast, vacuum melted HY130 heat C. Hot isostatic
pressing was effective at reducing the volume fraction of
porosity in heat C from 1.8×10-4 to 1.0×10-5 and reducing
the average pore or void size by 46%, from 2.7 to 1.4 µm.
Upon healing the large voids by HIP, the toughness
increased by 42% going from 114 kJ/m2 to 162 kJ/m2.
of the high strength steels in the current investigation was
a strong function of the density, type, morphology and
distribution of the nonmetallic inclusions. Cr and Mo
steels that were treated with Ti showed as much as a 100
kJ/m2 reduction in toughness in comparison with the Al
and Ca treated steels. The reduction in toughness
observed in the ferro-titanium treated steels was attributed
to void nucleation at coarse TiN particles, which fractured
during dynamic loading as well intergranular fracture
facilitated by eutectic Type II MnS inclusions. Addition
of rare earths in the form of misch metal increased the
toughness of the ferro-titanium heats to 88 kJ/m2 by
converting Type II MnS inclusions to globular MnS.
Deoxidation with Al and Ca produced mainly globular
calcium aluminate and Type I MnS inclusions. which
promoted ductile failure by void nucleation resulting from
decohesion of the particle from the matrix. Increasing the
inclusion spacing and average inclusion diameter
generally increased the toughness, while increasing the
density and volume fraction of inclusions decreased the
toughness. At equivalent inclusion diameter and
distribution, the Al and Ca deoxidized heats had much
higher toughness than the heats treated with
ferro-titanium. Addition of Ni from 1.6% in the 4325
steel to 5.4% in the HY130 resulted in a 10 kJ/m2 increase
in toughness in a steel with three times the inclusion
density. Therefore, for maximum toughness in cast steels
with higher inclusion contents, the Ni contents in Q and T
Cr and Mo steel should be kept at levels greater than 5%.
However, the benefit of additional Ni content was not
apparent for the steels treated with Ti. Therefore, the use
of Ti in Cr and Mo steels is strongly discouraged if high
toughness is required. However, if Ti treatment is
necessary for control of nitrogen, it should be
accompanied by additions of rare earth elements in the
form of misch metal to eliminate the formation of eutectic
Type II sulfides. The best practice is to use Al and Ca
deoxidation in heats prepared under a protective
atmosphere and cast using an appropriate filter.
CONCLUSIONS
ACKNOWLEDGMENTS
The dynamic fracture toughness of quenched and
tempered Cr and Mo steels with Ni contents of 0, 1.56
and 5.5 wt.% was evaluated with regard to deoxidation
practice and compared with the toughness values obtained
for a lightweight Fe-30.40%Mn-8.83%Al-1.07%Si0.90%C-0.53%Mo steel aged to an equivalent hardness.
The highest toughness, 165 kJ/m2, was obtained for the
5.5 wt.% HY130, which was Al killed, Ca treated and cast
utilizing a 64 ppi filter. A high amount of porosity
drastically decreased the toughness of the HY130 heat C
to 114 kJ/m2. However, hot istostatically pressing of
specimens from HY130 heat C was effective at closing
large voids and restoring toughness to 162 kJ/m2.
Statistically speaking, the 4325 and the lightweight
Fe-Mn-Al-C achieved equivalent dynamic fracture
toughness with values of 153 kJ/m2. At a lower hardness
of 32 HRC, the Fe-Mn-Al-C alloy obtained a much higher
toughness of 376 kJ/m.2 The dynamic fracture toughness
This work was supported in part by Army Research
Laboratory under contracts from Battelle Memorial
Institute (contract W911NF-07-D-0001) and Benet
Laboratories (contract W15QKN-07-2-0004). Laura
Bartlett was also supported by a U.S. Department of
Education GAANN fellowship under contract
P200A0900048. The following foundries are gratefully
acknowledged for providing test materials: Waukesha
Foundry, Inc., Nova Precision Casting Corporation and
Precision Castings, Inc. The authors also gratefully
acknowledge Bodycote Inc. for HIP of selected test bars
as well as Mr. Brandon Ensor, Mr. Tyler Preall, and Mr.
Joseph Brookshire for help with specimen preparation.
Paper 12-054.pdf, Page 16 of 17
AFS Proceedings 2012 © American Foundry Society, Schaumburg, IL USA
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