296
BEARING STEEL TECHNOLOGIES: 11TH VOLUME
STP 1600, 2017 / available online at www.astm.org / doi: 10.1520/STP160020160160
Mohamed Y. Sherif,1 Urszula Sachadel,1 Aidan Kerrigan,1
Boris Minov,1 Hanzheng Huang,1 Ilona Paape,1 and
Rene Gerritzen1
Novel Tough Micro-Alloyed
Bearing Steel with High
Hardenability
Citation
Sherif, M. Y., Sachadel, U., Kerrigan, A., Minov, B., Huang, H., Paape, I., and Gerritzen, R., “Novel
Tough Micro-Alloyed Bearing Steel with High Hardenability,” Bearing Steel Technologies: 11th
Volume, Progress in Steel Technologies and Bearing Steel Quality Assurance, ASTM STP1600,
J. M. Beswick, Ed., ASTM International, West Conshohocken, PA, 2017, pp. 296–322, http://
dx.doi.org/10.1520/STP1600201601602
ABSTRACT
A near-eutectoid lower-carbon vanadium-alloyed steel has been under
consideration for bearing steel components where improved toughness, and
therefore better micro-defect tolerance, is required. Due to the steel’s relatively
lower concentration of carbon—that is, compared with a typical approximately
1 wt.% C through-hardened bearing steel, the hardenability was found to be
higher. The steel design and heat treatment concepts have been investigated
experimentally using, for example, dilatometry and transmission electron
microscopy characterizing small-scale lab melts. The novel steel has been
designed with a particular focus on continuous casting as a suitable process.
Keywords
hardenability, micro-alloying, vanadium, continuous casting, bainite, transmission electron microscopy (TEM)
Manuscript received November 30, 2016; accepted for publication March 18, 2017.
1
SKF B.V., SKF Engineering & Research Centre, Kelvinbaan 16, 3439 MT Nieuwegein, The Netherlands
2
ASTM 11th International Symposium on Bearing Steel Technologies: Progress in Bearing Steel Technologies
and Bearing Steel Quality Assurance, on November 16–18, 2016 in Orlando, FL.
C 2017 by ASTM International, 100 Barr Harbor Drive, PO Box C700, West Conshohocken, PA 19428-2959.
Copyright V
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SHERIF ET AL., DOI: 10.1520/STP160020160160
Introduction
Candidate near-eutectoid bearing steels with a carbon content that is lower than,
for example, the well-known 100Cr6 bearing steel, could be beneficial to utilize
from a fatigue-resistance standpoint. The theory is that the small approximately 1
lm-sized carbides, which most probably are cementite particles embedded in the
hardened martensitic-tempered or bainitic matrix, could represent small internal
notches or discontinuities [1]. Consequently, Berns and Theisen [1] argued for the
use of near-eutectoid bearing steels over their regular higher carbon hypereutectoid
counterparts.
Meanwhile, for example, Bhadeshia and Solano-Alvarez [2] suggested increasing the bearing steel’s chromium content to stabilize the cementite phase thermodynamically. Therefore, perhaps it is beneficial to focus on at least the partial
replacement of cementite by the more thermodynamically stable types MC and
M7C3, where M stands for the metal content of the carbide.
A lower alloy carbon content ensures that more alloying elements, such as
chromium and molybdenum, are in solid solution in austenite during austenitization and are not tied in cementite. This leads to increased hardenability without the
need for raising the alloy content of such elements.
Nevertheless, there must be a minimum limit for the alloy carbon content
as the steel, for example, once quenched and tempered has to have a microstructure with sufficient hardness for adequate rolling contact fatigue performance. For example, a carbon content of about 0.5 to 0.6 wt.% is required in
solid solution in the austenite phase to ensure a hardness of more than 58
HRC [3]. The current design target hardness on bainitically hardened components is 60.5 HRC.
If transformation of the bearing components into bainite is desired, an additional alloy design constraint emerges, which is to ensure that the martensitestart temperature (MS) of the austenite matrix is sufficiently depressed below
the quenching temperature. As will be seen in the following sections, such a requirement was not easy to meet as it must be balanced with the hardenability
requirement mentioned earlier. Consequently, we investigated the option to alloy the steel with microalloying elements such as vanadium [4]. Microalloying
with vanadium allowed for relatively higher austenitization temperatures to be
used to ensure the depression of the MS temperature below the temperature at
which bainite starts to form without excessive austenite grain growth. Also,
austenitizing at higher temperatures increases the hardenability, which is a
crucial requirement for the manufacture of thick-walled large-size bearing
components.
The higher hardenability of the newly developed steel allows for the potential
replacement of several well-established bearing steels such as 100CrMo7-3,
100CrMo7-4, and 100CrMnMoSi8-4-6 targeting medium to large-size bearings
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(d being inner ring bore diameter greater than 120 mm). With the manufacturing
process in mind, the new steel had to be designed to suit the continuous casting
process in addition to the traditional ingot-casting process route. In other words,
even with such a high steel hardenability requirement, the segregation of carbon
had to be within acceptable limits across the transverse bar section in the ashomogenized and hot-rolled condition.
Experimental Procedures
All metallographic specimens were hot-mounted in resin followed by grinding
and then finish-polished using 1-lm diamond paste according to ASTM E3-11.
A Struers AbraPol-20 automatic preparation machine was used for this purpose.
The as-quenched untempered specimens had to be cold-mounted to avoid any
potential microstructural changes caused by heating. All metallographic specimens were subsequently cleaned and dried as per standard well-known practices,
then immediately etched. Unless otherwise stated, the etchant was Nital with a
concentration of 1.5 % (by volume).
Transmission electron microscopy (TEM) work was carried out on an FEI
Titan 200 kV microscope. TEM foils were initially thinned by means of grinding
and, finally, a Struers TenuPol twin jet electropolisher was used.
The X-ray specimens were prepared according to standard practices followed
for the preparation of metallographic specimens, with the addition of an electropolishing step carried out after mechanical polishing. Zr-filtered Mo Ka1,2 radiation
was used in a Bruker D8 DISCOVER diffractometer operated at 50 kV and 50 mA.
A step-scan mode was selected over the scanned 2h range of 26.58 to 408. The angular step width was 0.028 with a collecting time of 10 s at each step. The diffraction
lines a(200), c(220), a(211), and c(311) were considered for the determination of
the retained austenite content.
A push-rod TA Instruments DIL805A/D dilatometer was used in the quenching mode where only the temperature of the tested specimen was allowed to vary.
All dilatometer specimens were solid cylinders measuring 4 mm in diameter and
7 mm in length.
EXPERIMENTAL STEEL MELTS
As can be seen in Table 1, seven experimental steel melts have been investigated during the exploratory phase of the work. The experimental steel melts have all been
vacuum induction melted. Steels A, B, F, and G were processed such that the molten steel was cast into a mold that was removed from the vacuum furnace at around
6008C, then air-cooled, and subsequently stripped when cold. Per steel, two ingots
were prepared, approximately 30 kg each. The top and bottom of the ingots were
cropped and discarded. Homogenization was carried out in an air furnace at
1,2008C for 24 h. Afterward, the steels were furnace-cooled down to 3508C and
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TABLE 1 The chemical composition of the exploratory experimental steel melts and the
reference steel H.
Element,
wt.%
C
D
E
G
H (100CrMo73/IC)
0.75
0.79
0.97
0.34
0.34
0.26
0.81
0.66
1.79
A
B
F
C
0.73
0.73
0.72
0.78
0.83
Si
0.20
0.21
0.33
0.36
0.34
Mn
0.81
0.82
0.82
0.79
0.80
0.81
Cr
1.63
1.68
1.66
1.71
1.69
1.68
1.68
Ni
<0.001
<0.001
0.09
0.10
0.10
0.10
0.10
0.11
Mo
0.36
0.36
0.34
0.34
0.36
0.35
0.35
0.26
Cu
0.006
0.005
0.196
0.201
0.201
0.193
0.192
0.206
V
0.106
0.003
0.10
0.10
0.10
0.10
0.10
0.009
P
0.007
0.006
0.004
0.004
0.004
0.006
0.006
0.006
S
0.01
0.01
0.001
0.001
0.001
0.004
0.004
0.004
Al
0.006
0.033
0.003
0.003
0.002
0.012
0.012
0.028
As
0.001
0.001
0.003
0.003
0.003
0.001
0.001
0.015
Sn
0.001
0.001
0.001
0.001
0.001
0.001
0.001
0.011
Sb
<0.0015
<0.0015
0.0003
0.0004
0.0005
0.0001
0.0003
0.0025
Tia
15
15
3
3
3
5
5
16
Ba
4
4
2
3
2
3
3
2
Pba
<5
<5
2
2
3
1
2
<5
Caa
1
1
3
3
3
2
2
2
Na
118
79
150
150
133
134
118
50
Oa
8.3
6.0
30
16
19.5
9.4
10.6
3.7
Mo/Si
1.80
1.71
1.03
0.94
1.06
1.03
1.03
1.00
Cr/C
2.23
2.30
2.31
2.19
2.04
2.24
2.13
1.85
ICD, mm
123.6
125.7
131.4
130.8
131.2
133.5
132.3
92.5
Note: aIn ppm; IC ¼ ingot cast; C ¼ carbon; Si ¼ silicon; Mn ¼ manganese; Cr ¼ chromium;
Ni ¼ nickel; Mo ¼ molybdenum; Cu ¼ copper; V ¼ vanadium; P ¼ phosphorus; S ¼ sulfur;
Al ¼ aluminum; As ¼ arsenic; Sn ¼ tin; Sb ¼ antimony; Ti ¼ titanium; B ¼ boron; Pb ¼ lead;
Ca ¼ calcium; N ¼ nitrogen; O ¼ oxygen; Mo/Si ¼ molybdenum/silicon; Cr/C ¼ chromium/carbon;
ICD ¼ ideal critical diameter [5].
then air-cooled. The hot-forging process was carried out at a temperature within
the range 1,1008C to 1,2008C. The minimum forging temperature was 9008C followed by vermiculite cooling to ensure sufficiently slow cooling for the formation
of fine vanadium-rich precipitates. The forged bars were about Ø 30 mm with the
area reduction ratio of about 10:1.
The steels with the codes C, D, and E were sourced to carry out a sensitivity
analysis to determine the most suitable carbon range for the new steel concept.
The casting was executed by pouring the liquid steel into an adjacent mold placed
inside the vacuum chamber of the melting furnace. After cooling, in vacuum, the
removal of the cast head yielded 55 kg to 60 kg per ingot. Six blocks, approximately 125 mm by 125 mm by 80 mm in size, were sectioned per ingot. The blocks
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were then reheated at 1,2008C and then hot-rolled to a final thickness of approximately 20 mm. After hot-rolling, the plates were air-cooled to 6508C before being
allowed to continue to cool in a furnace that was kept at this temperature and
then switched off.
The reference Steel H (100CrMo7-3) was commercially produced by ingot
casting.
CONTINUOUSLY CAST STEELS
Two steel types were manufactured in cooperation with Deutsche Edelstahlwerke
(DEW), Witten, Germany, using an industrial vertical continuous caster. The chemical
compositions are shown in Table 2. Prior to hot-rolling, the steels were homogenized at
1,2788C for 6 h, for a total furnace time—including heating to temperature—of 12 h.
A variant of the new continuously cast steel (I) was homogenized for only 5.5 h of total
furnace time (3 h soaking at 1,2788C); however, it was later abandoned due to nonconformance with regard to segregation. Hot-rolling of steels I and J was carried out,
resulting in various bar diameters.
TABLE 2 The chemical composition of the continuously cast steels.
Element, wt.%
I (New Development)
J (100CrMo7-3)
C
0.80
0.94
Si
0.36
0.27
Mn
0.79
0.63
Cr
1.55
1.77
Ni
0.12
0.10
Mo
0.37
0.26
Cu
0.09
0.07
V
0.09
-
P
0.011
0.012
S
0.001
0.003
Al
0.026
0.009
As
0.004
0.004
Sn
0.004
0.003
Sb
< 0.0050
< 0.0020
Tia
13
13
Pba
< 10
< 20
Caa
<5
<5
Ha (liquid)
2.0
2.0
Oa
9.0
9.0
Supplier
DEW
DEW
Heat no.
586200
579030
a
Note: In ppm. H ¼ hydrogen.
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SHERIF ET AL., DOI: 10.1520/STP160020160160
Steel Design
The new steel chemical composition is based mainly on that of steel 100CrMo7-3
but with a lower carbon content and with the addition of vanadium (obviously
while attempting to meet all design requirements as mentioned in the introduction section).
It is known that the eutectoid carbon content is at about 0.8 wt.%, that is, for
the iron-carbon system. As can be seen in Fig. 1, as the steel is alloyed further, its eutectoid carbon content shifts to lower carbon concentrations. So even though the
carbon content of the newly developed steel was 0.8 wt.%, the steel is a hypereutectoid steel because its eutectoid carbon content is 0.72 wt.% (Fig. 2). The equilibrium
phase diagram of the reference steel whose chemical composition is shown in Table 2
is presented in Fig. 3.
For better clarity, Fig. 4 and Fig. 5 are reproductions of the equilibrium phase
diagrams for both steels with narrower carbon and temperature ranges.
As can be seen in Fig. 4, for the new steel with its carbon content of 0.8 wt.%,
other types of carbides such as VC (MC) and M7C3 (Cr-rich) can be present at
equilibrium along with cementite (M3C). At typical austenitization temperatures,
the cementite phase, at equilibrium, is completely unstable and is replaced with
the aforementioned types of carbides. The MC-type is known to be rather stable,
and when present as very fine precipitates, it prevents excessive austenite grain
growth [4].
The reference steel is strikingly different in that cementite is stable over the
temperature range shown in Fig. 5.
FIG. 1 The effect of alloying elements on the eutectoid carbon content [6]. (Source:
Republished with permission of ASM International, from ASM Handbook, Vol. 4,
Torsten Ericsson, 1991; permission conveyed through Copyright Clearance
Center, Inc.)
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FIG. 2 The equilibrium phase diagram of the newly developed Steel I. The arrow points
to the eutectoid carbon content of 0.72 wt.%.
DESIGN FOR REDUCED SEGREGATION
Table 1 shows two ratios. The first is the Mo/Si ratio, which should be maximized to
reduce the formation of segregation channels during solidification [7]. The second
is the Cr/C ratio, which should also be high to ensure lower chromium segregation
during solidification (Fig. 6).
As can be seen in Fig. 6, the measured versus calculated chromium segregation
ratios for a given steel with a fixed chromium content of 1.5 wt.% are presented. It
can be seen in the figure that, first, different cooling rates—which affect the fraction of residual liquid (fliquid)—shift the peak’s position above which the chromium segregation ratio starts to decrease. In all cases, the chromium segregation
ratio does seem to peak at relatively high carbon contents. So, for a fixed chromium content, lower carbon results in less microsegregation. The calculations in
Fig. 6 suggest a decrease in the chromium segregation ratio at relatively high carbon contents [8]. This could be attributed to the formation of the eutectic, which
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 3 The equilibrium phase diagram of the reference Steel J. The arrow points to the
eutectoid carbon content of 0.70 wt.%.
dilutes the rejected chromium, leading to a lower concentration of the element in
the residual liquid [8].
As steels with carbon contents greater than 0.51 wt.%—according to the ironcarbon (Fe-C) equilibrium phase diagram—are cooled, the only solidification reaction that occurs is the formation of the austenitic dendrites [9]. Macrosegregation is
known to take place when the interdendritic liquid steel is absorbed into cavities
[10]. As such, the fluidity of the residual liquid will control the severity of the macrosegregation. Indeed, Takahashi, Kudoh, and Nagai, in their work on the solidification of carbon steel reported lower fluidity in a silicon-free 1 wt.% molybdenum
steel [11], which seems to support maximizing the Mo/Si ratio.
Also, lowering the carbon content, as is the case with the concept steels, would
enable the reduction of the solute-enriched liquid fluidity and thus reducing macrosegregation, rendering the steels more suitable for continuous casting compared
with conventional higher carbon bearing steels. Oh and Chang evaluated liquid
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FIG. 4 The equilibrium phase diagram of the newly developed Steel I.
steel fluidity in the mushy zone in terms of specific permeability (k) and found that
its calculated value decreased as the steel carbon content decreased [10].
Discussion
AS-RECEIVED HOT-ROLLED STRUCTURES
As expected, all the steels exhibited a fully pearlitic microstructure; see, for example,
Fig. 7. The as-received hardness of the plates can be seen in Fig. 8. As can be seen in
the figure, within experimental error, it can be concluded that there is not a significant difference in hardness among these three steels in the hot-rolled pearlitic
condition.
SPHEROIDIZING-ANNEALING PROCESS
Full annealing was carried out to ensure soft-machinability and a better response to
austenitization during hardening. Using the steels shown in Fig. 7, whose carbon
contents were 0.72 wt.%, 0.78 wt.%, and 0.83 wt.%, it was possible to discern how
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FIG. 5 The equilibrium phase diagram of the reference Steel J.
robust the steel design is with regard to the spheroidizing-annealing process. For
this purpose, a spheroidizing-annealing schedule typical for bearing steels was used
[12] but was adapted for the current chemical compositions, mainly by varying the
annealing temperature. Therefore, it was possible to investigate the effect of the alloy carbon content versus the annealing temperature.
Fig. 9 presents the effect of varying the annealing temperature on the measured “soft” hardness for the steels indicated. For ease of soft-machinability, the
acceptable hardness range was set to be 180 to 220 HV. As such, the hardness of
all the annealed microstructures was deemed acceptable, with the difference in
alloy carbon content combined with the variation of the annealing temperature
not causing any nonconformance. Only carbon was varied in the sensitivity
to spheroidizing-annealing work due to that element being the fastest diffusing.
Fig. 10 shows the full annealed microstructures obtained when spheroidizingannealing Steel D as an example.
Additionally, Steels C, D, and E were quenched to form martensite and tempered, from the various spheroidized conditions by austenitizing at 9058C for
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FIG. 6 Alloy carbon content (with Cr ¼ 1.5 wt.%) versus the calculated chromium
segregation ratio (XCr,max/XCr,min) for the cases shown with various percentages
of residual liquid. The data points are experimental results from references
4 and 5 in Kozeschnik [8]. (Source: Metallurgical and Materials Transactions A,
“A Scheil-Gulliver Model with Back-Diffusion Applied to the Microsegregation
of Chromium in Fe-Cr-C Alloys,” Vol. 31, No. 6, 2000, Ernst Kozeschnik, with
permission of Springer.)
70 min in a high-temperature salt bath furnace, oil quenched, and then cooled to
room temperature.
The hardened coupons were about 18 mm thick. Martensitic hardening was
chosen in this case because it is fast compared with transformation into bainite.
Half of the specimens were left in the as-quenched condition, whereas the other
half was immediately tempered by heating to 2008C and kept at temperature for 2 h
before air-cooling. The implemented heat treatment schedule was a variation from
that typically used to heat-treat 100Cr6 bearing steel components, while taking into
account the outcome of prior trials. The objective was to investigate the effect of the
variation in the alloy carbon content and the adopted spheroidizing-annealing
cycle—not only on the microstructure in the soft condition as mentioned earlier
but in the as-quenched and quenched and tempered conditions as well.
As shown in Fig. 11, in the as-tempered condition, Steel C (0.72 wt.% carbon)
showed the lowest average hardness of 70161 HV10, while the variants with 0.78
and 0.83 wt.% carbon exhibited a higher average hardness of 72064 and 72563
HV10, respectively. For all the hardened steel structures studied, it can be concluded that the hardness measured was insensitive to the annealing temperature
used.
Table 3 presents the retained austenite measurements made on the annealed
steels in both the quenched and quenched and tempered conditions. The
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 7 The as-received hot-rolled pearlitic microstructures of Steels (a) C (b) D, and (c)
E. The micrographs were taken from the mid-thickness of the plates. The rolling
direction is horizontal. The measured hardness is shown in Fig. 8.
FIG. 8 The measured hardness of the as-received steel microstructures shown in Fig. 7.
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FIG. 9 The measured hardness of the fully annealed microstructures.
annealing temperature of 8108C was selected in the study as it exhibited the best
microstructure in terms of minimum lamellar pearlite content and optimum carbide size. Given the relatively low carbon in the new steel concept, undissolved
grain boundary carbides were never an issue as their diameter was typically
much smaller than 12 lm. The somewhat high retained austenite contents shown
in Table 3 can be attributed to the elevated austenitizing temperature of 9058C—
an explanation for the usage of such a temperature will follow. Nevertheless, as
may be expected, as the carbon content of the steel increased, the retained austenite content should increase because the MS temperature is expected to be
depressed.
OPTIMIZATION OF THE AUSTENITIZATION CONDITIONS
An essential aspect of the new alloy concept is microalloying with vanadium. In the
early phase of the work on Steels A and B, given the objective to transform the steels
into bainite, it was crucial to ensure that the MS temperature was depressed below
the bainite transformation temperature. With the carbon content of the steels being
significantly lower than the reference Steel H (Table 1), this requirement was difficult
to meet unless higher austenitization temperatures were used in the hardening heat
treatment. When austenitizing Steel A at 9008C for 50 min, the dilatometermeasured MS temperature was about 2008C, while Steel B austenitized at 8958C for
50 min had an MS temperature of about 2038C. This is in contrast to the reference
Steel H that was austenitized at a lower temperature of 8658C for 50 min and exhibited an average MS temperature of 2088C. Most probably the lack of copper and
nickel residuals in these two experimental melts, both of which are austenite stabilizing elements, compared with the commercial 100CrMo7-3 steel (H) caused the
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 10 Visible light optical microscopy images of the microstructure of Steel D
annealed at the indicated temperatures. The measured hardness values can
be seen in Fig. 9.
observed rise of the MS temperature. The lack of residuals prompted the present
authors to order the manufacture of the last series of experimental melts coded F
and G (Table 1).
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FIG. 11 The measured hardness in the as-quenched and tempered conditions. See text
for details on the corresponding heat treatment schedules. The applied prior
annealing temperatures are indicated.
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TABLE 3 The X ray measured retained austenite content, in volume percent, of the indicated steels
fully annealed at 8108C.
Condition/Steel
C
D
E
As-quenched
16
19
21
Quenched and tempered
14
16
21
Nevertheless, the vanadium-alloyed Steel A showed finer prior austenite grain
sizes compared with its vanadium-free counterpart and the reference steel even
while soaked at the highest austenitization temperature (Fig. 12 and Fig. 13).
VANADIUM-RICH PRECIPITATES
Fig. 14 shows a bright field image of the bainitically hardened Steel A component
with vanadium-rich precipitates visible (circled). The energy dispersive X-ray spectroscopy spectra are presented in Fig. 15. The particles marked “1” and “2” were confirmed to be vanadium carbides. The vanadium carbides were very fine, as
expected, which could explain the small prior austenite grain sizes obtained for the
same steel (Fig. 12).
FIG. 12 The prior austenite grain sizes (PAGS) of the quenched and tempered steel
specimens. The microstructures soaked for 70 min at temperature are
presented in Fig. 13.
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FIG. 13 The prior austenite grains of the tempered martensitic microstructures of
Steels (a) A austenitized at 9058C, (b) B austenitized at 8958C, and (c) H
austenitized at 8958C, all soaked for 70 min at temperature.
CONTINUOUSLY CAST INDUSTRIAL MELTS
As mentioned earlier and shown in Table 2, two steel melts have been ordered with
the chemical composition of the new steel concept and the reference steel. For the
following work, 120-mm diameter bars were used.
Both melts were produced in a 130-ton electric arc furnace using presorted steel
scrap. The standard secondary refining process was composed of ladle treatment
and vacuum degassing. The casting took place in a vertical continuous caster via
four nozzles inserted into water-cooled open-ended copper molds that oscillated
along the axis of the moving strand of 475 mm by 340 mm cross section. The solidified strands were cut into blooms. The cold blooms coming from the steady-state
casting conditions were subsequently homogenized, as described previously, and
then hot-rolled into the required bar diameters. Although not presented here, the
quality assurance including, for example, chemical composition analyses, micrononmetallic inclusion characterization, macrosegregation, and carbide network
assessments were carried out on 120-mm diameter bars, for both melts, and were
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FIG. 14 Bright field transmission electron microscopy image of Steel A bainitic
microstructure with vanadium-rich precipitates circled.
found to be in conformance with SKF’s internal specifications for hot-rolled bars
for high-carbon, through-hardenable bearing steels.
MECHANICAL PROPERTIES
After spheroidizing-annealing of the bar material, the Steel I soft-machined test
components were transformed into bainite (bainitic-ferrite) according to four different schedules named Bainite 1 through Bainite 4.
As a comparison, test components from Steel J were also bainitically heattreated; there was only one variant, which was processed in accordance with SKF’s
standard practices for this steel type. The impact toughness test results are presented
in Fig. 16, where it can be seen that the Bainite 2 variant of the new steel showed
increased impact toughness (average 19 J) over the reference steel microstructure
(average 13 J). These results are reproduced versus the measured hardness—in addition to results from past test campaigns for comparison, as shown in Fig. 17. From
Fig. 17, it is shown that in terms of hardness and absorbed energy balance, the
Bainite 2 heat treatment schedule had the best combination of these two properties.
Consequently, although not stated here, the bainite transformation schedule
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FIG. 15 The qualitative chemical analyses of the precipitates shown in Fig. 14.
“Bainite 2” was adopted. This bainite through-hardening heat treatment process
was also the most suitable for series production.
HARDENABILITY
Dilatometry
In this test, the steel concept was investigated against not the reference steel but a
higher alloyed steel grade that has higher hardenability, the 100CrMo7-4 steel.
For this purpose, bar material from the latter material was obtained from Ovako,
Sweden (heat no. D4696), which had the chemical composition Fe-0.95C-0.31Si0.64Mn-1.71Cr-0.23Ni-0.41Mo-0.191Cu (wt.%). The bar material was given the
code “K.”
From the fully annealed condition, dilatometer specimens were machined from
both steels in the axial direction and from the middle of the bar radius. As can be
seen in Table 4, two austenitization temperatures were chosen, combined with several t8/5 times. The t8/5 times correspond to certain bar diameters quenched in oil
[13,14].
Two specimens were investigated per steel per test condition. The results’ reproducibility was very good, and testing more than two specimens per condition
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 16 The absorbed energy of the bainitically hardened steel specimens.
FIG. 17 The blunt-notch impact toughness test results versus measured hardness. The
arrows point to the properties of the continuously cast industrial steels whose
chemical composition is shown in Table 2.
was deemed unnecessary. Given the high hardenability of the steels, only the structures austenitized at 8908C will be discussed hereafter. Also, as indicated in Table 4,
for certain austenitization conditions, some t8/5 times were considered to be either
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TABLE 4 The hardenability test matrix executed on the dilatometer for both the steel concept
(Steel I) and a 100CrMo7-4 melt (Steel K); t8/5 is the cooling time between 8008C and
5008C, in seconds.
Bar Diameter, mm/Oil
[13,14]
t8/5, s
8708C/70 min
8908C/70 min
25
35
H
NA
50
60
H
NA
70
75
H
NA
125
110
H
NA
200
130
H
H
300
170
H
H
500
220
NA
H
too long or too short. Fig. 18 shows an example of the obtained results where it was
clear that the two steels had hardenability characteristics different from the
100CrMo7-4 steel specimens, transforming at higher temperatures, upon cooling,
at approximately 4508C into products of very fine pearlite and probably upper bainite, all of which are undesirable phases to obtain in the as-quenched microstructure
FIG. 18 The dilatometer curves obtained by austenitizing at 8908C for 70 min followed
by cooling with t8/5 value of 500 s. The relative change in length equals strain
times 100.
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SHERIF ET AL., DOI: 10.1520/STP160020160160
(Figs. 19–21). All specimens had been austenitized at 8908C for 70 min followed by
cooling with t8/5 of 500 s (Fig. 18). The concept steel shows a deviation from linearity upon cooling, indicating transformation. However, this indication of transformation was difficult to observe on the dilatometer curves under the same
magnification.
All the microstructures shown in Figs. 19 through 21 were observed in the center
of longitudinally sectioned specimens. In light of these results, it was necessary to
carry out Jominy testing with the inclusion of an established bearing steel material
with a higher hardenability, such as the 100CrMnMoSi8-4-6.
End-Quench (Jominy Test)
The hardenability of the new steel concept has been assessed according to the ISO
642:1999 test specification using specimens machined from normalized 120-mm diameter bars. In addition to the steels shown in Table 2, the steels in Table 5 were also
included in the Jominy test campaign.
The Jominy test results are presented in Fig. 22. Four tests were carried out
quenching the test rods from the temperatures 8658C, 8708C, 8908C, and 8958C.
FIG. 19 Visible light optical microscopy images showing the dilatometer-quenched
microstructures after austenitizing at 8908C for 70 min followed by cooling
with t8/5 value of 200 s: (a) and (b) Specimens 1 and 2 from Steel I,
respectively; (c) and (d) Specimens 1 and 2 from Steel K, respectively.
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FIG. 20 Visible light optical microscopy images showing the dilatometer-quenched
microstructures after austenitizing at 8908C for 70 min followed by cooling
with t8/5 value of 300 s: (a) and (b) Specimens 1 and 2 from Steel I,
respectively; (c) and (d) Specimens 1 and 2 from Steel K, respectively.
The Steel J (100CrMo7-3) was quenched from all the temperatures except
8958C as it was considered rather high for this alloy system. The opposite was true
for Steels L (100CrMnMoSi8-4-6) and M (100CrMo7-4), where none of the steel
rods was austenitized at 8658C, which was deemed too low for such relatively highalloyed grades. This test temperature is believed to be quite low for the new steel
(I); nevertheless, the steel was included for comparison.
As presented in Fig. 22a, Steel I clearly outperforms the reference Steel J, both
quenched from 8658C. There is a significant increase in hardness of the quenched
new steel near the head of the rod, far from the quenched-end when austenitized at
8708C, which was not the case with the reference steel (Fig. 22b). Steel M exhibited a
somewhat higher level of hardenability compared with the reference steel (J), which
was expected. The relatively high-alloyed Steel L showed the highest hardness measured over approximately the other half of the rod away from the quenched end. By
austenitizing at 8908C (Fig. 22c), all the steels had the same ranking as in Fig. 22b;
nevertheless, it is clear that the novel steel austenitized at this temperature exhibited
hardenability similar to that of the most highly alloyed steel studied herein—Steel L
(100CrMnMoSi8-4-6).
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 21 Visible light optical microscopy images showing the dilatometer-quenched
microstructures after austenitizing at 8908C for 70 min followed by cooling
with t8/5 value of 500 s: (a) and (b) Specimens 1 and 2 from Steel I,
respectively; (c) and (d) Specimens 1 and 2 from Steel K, respectively.
Upon austenitizing at an even higher temperature (8958C), the new steel
appeared to possess hardenability characteristics similar to the 100CrMnMoSi8-4-6
steel (Fig. 22d).
So, Steel I exhibited high hardenability, at least on a par with the
100CrMnMoSi8-4-6 steel. Perhaps an observation would be that the new steel
appears to show continuously improving hardenability. In other words, in this
work, unlike the case for the 100CrMnMoSi8-4-6 steel, the new steel showed a difference in hardenability if the steel was austenitized at 8908C or 8958C. This perhaps
could be attributed to the cementite in the new steel microstructure being more easy
to dissolve compared with other more thermodynamically stable carbides such as
the MC-type, thereby making elements such as molybdenum and chromium more
available—in solid solution—to increase the hardenability of the austenitic matrix.
Nevertheless, this 58C difference in the test temperature is rather small and most
likely is well within the expected variation in temperature in industrial furnaces.
This hardenability performance of the new steel was positively surprising, and
given that it is not known to the present authors if the 100CrMnMoSi8-4-6 steel
can be continuously cast, then perhaps the current steel concept possesses a significant economic advantage over the former steel.
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TABLE 5 The chemical composition of two additional steels considered for the Jominy test.
Element, wt.%
Steel L (100CrMnMoSi8-4-6)
Steel M (100CrMo7-4)
C
0.94
0.96
Si
0.43
0.27
Mn
0.96
0.68
Cr
1.92
1.67
Ni
0.24
0.16
Mo
0.54
0.42
Cu
0.176
0.117
V
0.011
0.011
P
0.011
0.010
S
0.006
0.008
Al
0.019
0.021
As
0.009
0.007
Sn
0.013
0.008
Sb
0.0027
0.0029
Tia
15
14
Pba
1
7
Caa
3
3
Oa
6.1
Supplier
Ovako
Ovako
Heat no.
D9075
G3960
Note: aIn ppm.
Conclusions
The following points may be summarized:
1. A new through-hardenable steel has been designed for medium- to large-size
bearings (d being inner ring bore diameter greater than 120 mm) exhibiting
superior hardenability perhaps compared with that of the well-established
100CrMnMoSi8-4-6 steel. As such, the new steel could replace bearing steels
100CrMo7-3, 100CrMo7-4, and possibly at least partly the 100CrMnMoSi8-46 steel.
2. In addition to being castable via ingot-casting, the new steel is designed with
continuous castability in mind. So far, a continuous vertical caster has been
used in the manufacturing of the steel, and the possibility of utilizing a bowtype caster would need further investigation.
3. The necessary heat treatment schedules have all been developed, resulting in
better impact toughness in the bainitic-hardened condition, with good hardness, compared with the reference 100CrMo7-3 steel.
4. Given the new steel’s somewhat lower carbon content, it may offer the potential for component surface thermochemical treatments such as carbonitriding
and carburizing. The compatibility of the steel concept with these processes is
therefore recommended for future study.
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SHERIF ET AL., DOI: 10.1520/STP160020160160
FIG. 22 The Jominy hardenability curves obtained according to ISO 642:1999 for the
indicated steels following quenching from (a) 865 C; (b) 870 C; (c) 890 C; and
(d) 895 C.
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ACKNOWLEDGMENTS
The authors would like to acknowledge support from Dr. Steve Ooi, SKF UTC, University of Cambridge, in the TEM work.
References
[1]
Berns, H. and Theisen, W., Ferrous Materials, Springer-Verlag, Berlin, 2008.
[2]
Bhadeshia, H. K. D. H. and Solano-Alvarez, W., “Critical Assessment 13: Elimination of
White Etching Matter in Bearing Steels,” Mat. Sci. Tech., Vol. 31, No. 9, 2015, pp. 1011–1015.
[3]
Berns, H., Gavriljuk, V., and Shanina, B., “Intensive Interstitial Strengthening of Stainless
Steels,” Adv. Eng. Mater., Vol. 10, No. 12, 2008, pp. 1083–1093.
[4]
Gladman, T., The Physical Metallurgy of Microalloyed Steels, Institute of Materials,
London, 2002.
[5]
Jatczak, C. F., “Hardenability in High Carbon Steels,” Metall. Trans., Vol. 4, No. 10, 1973,
pp. 2267–2277.
[6]
ASM Handbook, Heat Treating,” Vol. 4, ASM International, Novelty, OH, 1991.
[7]
Lund, T., formerly of AB SKF, personal communication.
[8]
Kozeschnik, E., “A Scheil-Gulliver Model with Back-Diffusion Applied to the Microsegregation of Chromium in Fe-Cr-C Alloys,” Metall. Mater. Trans. A, Vol. 31, No. 6, 2000,
pp. 1682–1684.
[9]
Louhenkilpi, S., “Continuous Casting of Steel,” Treatise on Process Metallurgy, S.
Seetharaman, Ed., Elsevier, Boston, 2014, pp. 373–434.
[10] Oh, K. S. and Chang, Y. W., “Macrosegregation Behavior in Continuously Cast High Carbon Steel Blooms and Billets at the Final Stage of Solidification in Combination Stirring,”
ISIJ Int., Vol. 35, No. 7, 1995, pp. 866–875.
[11]
Takahashi, T., Kudoh, M., and Nagai, S. I., “Estimation of Effective and Specific Permeability in the Mushy Zone of Steel Ingot by the Method of Seepage into Cylindrical Hole,”
Tetsu-to-Hagané, Vol. 68, No. 6, 1982, pp. 623–632.
[12] Verhoeven, J. D., “The Role of the Divorced Eutectoid Transformation in the Spheroidisation of 52100 Steel,” Metall. Mater. Trans. A, Vol. 31, No. 10, 2000, pp. 2431–2438.
[13] Totten, G. E., Bates, C. E., and Clinton, N. A., Handbook of Quenchants and Quenching
Technology, ASM International, West Conshohocken, PA, 1995.
[14] Larsson, S., AB SKF, personal communication.
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