Effect of Mo in TRIP Steels (PR 311 075)

Prévia do material em texto

1.
2.
3.
The Effect of Mo in Si-Mn Nb
Bearing TRIP Steels
Michael Bouet], John Root2,
Elhachmi Es-Sadiqui3, Steve Yuel
McGill University, Dept. Of Metallurgical
Engineering, M.H. Wong Building,
Metallurgical Wing, 3610 University Street,
Montreal, Quebec, Canada, H3A 2B2
Neutron Program for Materials Research,
Chalk river Laboratories, Chalk River,
Ontario, Canada, KOJ 1JO
Canmet, 555 Booth Street, Ottawa, Ontario,
KIA OG1
det
ste
ste
tra
reta
pro
dra
tra
lev
as:
(i.e
inv
und
step
(TM
sph
mic
car
are obtained this way. Once cold formed, the product
undergoes a quench stage whereby it is heated up into the
austenite stable region, allowed to fully transform, and
quenched to generate a fi.dly martensitic structure. Lastly,
the product is tempered to provide the steel with a degree
of toughness. The cold forming of rod, intended for high
strength fasteners, is a classical example of a material
requiring these treatments.
Unlike heat treated steels, TMProcessed
Transformation Induced Plasticity (TRIP) steels offer the
advantage of generating the kind of formability required
from the steel in the as hot rolled condition, while
obtaining the mechanical properties required directly
from the cold forming operation. Hence, the stages of
spheroidizing, quenching and tempering would no longer
be required, thereby working towards satisfying the
incessant drive for high quality and cost effective new
classes of steels.
High levels of ductility by ‘TRIP’ occurs by the
Strain Induced Transformation (SIT) of metastable
Retained Austenite (RA) to the lower energy state
martensite phase. This transformation is accompanied by
a local increase in the strain-hardening rate, strengthening
the material and shifting the point of plastic instability
elsewhere and to higher strains. The fwst steels known to
T
ABSTR4CT
An experimental study was performed to
ermine the effect of Mo in Si-Mn Nb bearing TRIP
els, essentially as a partial replacement for Si. Three
el compositions were investigated. Bainite
nsformation conditions were investigated on the
ined austenite characteristics and final mechanical
perties. Results revealed that Mo has a strong solute
g effect, consequently retarding the ferrite and pearlite
nsformations. The Mo bearing steel (with lowered Si
el) generated excellent mechanical properties (as high
UTS=1269 MPa, T.El.= 35.9°/0)in the range of those
. no Mo, standard Si level) observed by previous
estigators.
1. INTRODUCTION
High strength C-Mn steel products typically
ergo a sequence of energy intensive heat treatment
s. Firstly, following ThermoMechanical Processing
P), prior to shaping, the steel undergoes a partial
eroidization stage whereby the as rolled ferrhdpearlite
rostructure is transformed to a ferrite/spheroidized
bide structure. Much improved formability properties
40
TRIP by the SIT of RA were identified by Zackay.[l)
Work done at McGill University has shown tha~
with controlled TMP of Si-Mn TRIP steels microalloyed
with Nb, impressively high tensile strengths (above 1200
MPa) can be achieved having a wide ductility range (20-
55%T.E1.).(24) Furthermore, optimum RA characteristics
and mechanical properties were achieved when processed
in the bainite transformation region, as was reported in
previous investigations (Figure 1).(5-8)
{
+A
1350 A+ A
20 25 M
~,E1. (%jj U W 5
Fig. 1. Range of mechanical properties observed for
Si-Mn TRIP steels by various investigators. Effect of
Nb addition on the mechanical properties of TRIP
steels.
H MWSP CONF. PROC., 1SS 1998 – 675
The level of stability of RA against the SIT has
a significant effect on the mechanical properties of TRIP
steels. Microstructure bearing RA having a wide stability
range will progressively transform as the material is
deformed, thereby progressively avoiding plastic
instability and generating greater strains to failure. Work
done on the stability of RA has shown that it is strongly
dependent on its particle size(z’s),neighboring phases(2’9),
chemical composition, and work hardening (i.e.
z1 IO) me c e~c~ent of residualdislocation density).( ‘ ‘
austenite during transformation has been the target of
many of these investigations, Many of these studies have
shown a strong correlation of C content on the retention
of austenite to room temperature as well as on its stability
against the SIT. The strong lattice distortion effect of this
interstitial element impedes the shearing nature of the
martensitic transformation. Si, in high levels ( 1.5-2.0
wt.VO) was reported to stabilize a significant amount of
RA.(’ ‘-]3)More recently, Tsukatani et al. have observed
similar trends where both the RA characteristics and
mechanical properties were optimized with steels
containing 2.0 Wt.O/OSi and 1.5 Wt.O/OMn.(lOJSi inhibits
cem
aus
exc
ten
det
pur
Wh
im
me
hav
the volubility of the carbonitride, which in turn is a
consequence of the net effect that Mo has on the activity
coefficients of Nb, C and N. These observations are
consistent with the work of Wada et. al.(lG),who showed
that Mo decrease the activity of C in austenite, which
effectively, lowers the driving force for precipitation.
Furthermore, the change in activity will affect the
diffisivity of the carbide forming elements. Elements
which increase the volubility of carbides or nitrides in
austenite will decrease the diffhsivity coefficients and
hence, the diffisivity, of the precipitating species.
With these fiidings in mind, it is suggested that
Mo, potentially, can be an important addition to TRIP
steels. Amongst the possible benefits include: (1) greater
austenite saturation levels with C and Nb as well as
increased effectiveness of these elements by the delayed
formation of cementite and Nb(C,N); (2) easy process
control through delays in transformation, and; (3) lower
Si levels required.
The effect of composition, in particular Mo and
Si additions, and TMP (i.e. bainite hold time and
temperature) on the RA characteristics, mechanical
s
o
S
t
t
p
t
u
c
c
r
m
M
1
1
1
precipitates and is, therefore, limited to small additions.
The focus of this work is to consider the
feasibility of replacing Si with Mo in the steel. A number
of investigators have reported that Mo in steel results in
significant recrystallization and precipitation delays~ 14”5)
The delays in recrystallization are attributed to a solute
effect. However, as the processing temperature is
decreased, the initiation of carbonitride precipitation will
accentuate this delay (i.e. when the recrystallization and
precipitation start curves intersect). Additionally, it has
been reported that Mo, when present, retards the
precipitation of Nb(CN). Mo is responsible for increasing
Table 1. Steel Chemical Co
Steel c Si Mo
A 0.18 1.66 ---
B 0.20 1.6 0.30
c 0.23 1.1 0.33
676 – 40TH MWSP CONF. PROC., 1SS 1998
rom differing combinations of Si and Mo. Hence, the
ffect of each of these constituents can be examined
eparately while keeping in mind the ultimate objective
f partially replacing Si with Mo in the steel. The level of
i in these steels is two to four times greater than those
ypical of plain cmbon steels. The Mo levels utilized, on
he other hand, are typical of industrial microalloyed steel
ractices. The steels were cast at Canmet and received in
he as hot rolled condition.
TMP and mechanical testing was performed
sing a materials Testing System (MTS) adapted for hot
ompression. ‘2’17)Samples were machined into small
ylinders 11.4 mm in height and a 1.5 heightidiameter
atio.
positions (wt.%)
n Nb Al N (ppm)
.4 0.025 0.053 40
.4 0.041 0.028 55
.6 0.036 0.036 40
entite formation allowing for greater degrees of
tenite saturation with C. On the hand, in levels
eeding 1.0 wt.’??o,it is responsible for an undesirable
acious oxidelayer yielding a poor surface quality and
eriorated mechanical properties if rolled into the steel.
Nb is generally added to the steel for the
pose of austenite pancaking and grain refinement.
en in solid solution, it was reported to significantly
prove the RA characteristics and, hence, the
chanical properties of the steel.(24) It does however
e a strong affinity for C and N to form Nb(C,N)
properties and microstructure of the steel are reported in
this work.
2. EXPERIMENTAL PROCEDURE
Three steels were investigated in this work
(Table 1). They each have similar compositions apart
f
e
2.1. Continuous cooling compression testing
The material behaviors were characterized by a
Continuous Cooling Compression (CCC) tectique(2,17,20),
whereby following solutionizing, the sample is cooled
(0.5°C/see) and strained (0.005/see) simultaneously, as is
shown in Figure 2. The simultaneous cooling and
deformation allow for pinpointing critical temperatures
where microstmctural variations occur in the steel. These
variations are visible as sharp inflections in the change in
flow stress of the material as it is cooled (i.e. stress versus
temperature curve).
2.2. Effect of bainite hold time and temperature
Figures 3 and 4 illustrate the generic TMP
schedule utilized for comparing the effects of bainite hold
temperature and time on the RA characteristics,
mechanical properties and microstructure of the three
steels. Following the solutionizing stage (i.e. 1200”C, 30
min.), the sample is cooled to 105O”C, a typical TMP
temperature, held for a couple of minutes and strained in
two steps by compression for grain refinement.
A
E,= s,= 0.3
g
650°C,-25 vol.%a
51X1°C,5min
3
5 3CO”C,5mi
Fig. 3. TMP schedule utilized for comparing the
effects of bainite hold temperature.
t c1= E*=0.3
E=O.1/seclo50”c
T
. ...... ............ ........ . . . . . . ........ ........ ... .,
3“clsec
... ............................ . ........ .. .................. ........ ....
Ar3
0T
Subsequently, it is cooled below the austenite-to-ferrite
transformation temperature (Arq) and held at 650”C to
obtain roughly 25 VOl.O/Oferrite. The hold duration is
dependent on the steel chemistry. Lastly, the sample is
quenched in a salt bath and isothermally held at various
temperatures (i.e. 300,400 and 500°C), for various times
(2, 5 and 10 min.), and then removed and air cooled.
AI= 0.5”Ckc
TIME
Fig. 2. Schematic illustration of the TMP schedule
utilized for the CCC Test.
4
k2
i ~m’
k
TIME
Fig. 4. TMP schedule utilized for comparing the
effects of isothermal bainite hold time.
2,3 Material Characterization
Microstructure were characterized using optical
microscopy. The RA level for each specimen was
quantified by neutron diffraction pattern analysis, using
the NRU reactor DUALSPEC powder diffractometer at
the Chalk River laboratories.( ‘8) The neutron beam
wavelength utilized was 0.13286 nm. The specimens
were rotated continuously within the beam to average out
any texture effect. The diffraction patterns were analyzed
by the Rietveld profile refinement method to extract
lattice parameters and relative volume fractions of the RA
H MWSP CONF, PROC., 1SS 1998-677
and ferrite phases in each specimen. The RA carbon
content was determined from its lattice parameter
measurement, aw, through the relation:
aw = 3.578nm +(),&%mx”/oCw
2.4 Mechanical properties
Room temperature mechanical properties were
determined by a shear punch technique. (19)This technique
allows for generating mechanical property measurements
from small sized specimens. A 3 mm diameter disk is
punched out of a thin slice of the sample (-350 pm
thick), while the MTS system monitors the
load/displacement behavior. The raw data are eventually
converted to equivalent tensile properties using a
calibration factor.
3. RESULTS AND DISCUSSION
3.1 Continuous cooling compression testing
to Mo as a solid solution strengthener, as well a solute
drag element which, consequently, retards dynamic
recovery. More importantly, however, Mo appears to
lower the beginning of pearlite formation temperature
(i.e. Arl) by roughly 35°C (this is characterized by a
sharp increase in the flow stress following the sudden dip
in the curve as the sample is cooled). This is in agreement
with other investigators(14- lG),who have shown that MO
has a strong retardation effect on both the kinetics and
thermodynamics of carbide formation in steel. In
addition, partial removal of Si (steel C) from steel B shifts
the initiation of cementite formation to even lower
temperatures. This is explained by the partial loss of some
of the ferrite stabilizing effect of Si.
3.2 Isothermal ferrite transformation
Isothermal ferrite transformation curves were
generated for each of the materials under investigation at
a temperature of 650”C (Figure 6). From this, the
respective ferrite soak time required for a ferrite volume
fraction, prior to quenching in the bainite transformation
region, could be determined. Because of the excessively
l
f
t
c
A
m
b
F
t
Continuous cooling compression testing was
performed for each of the above mentioned steels. Slope
inflections in the flow stress versus temperature curves
indicated the critical temperatures where variations in the
microstructure occur. Figure 5 illustrates and compares
the CCC curves for these three steels. It is observed that
the addition of Mo (steel B versus steel A) significantly
increases the flow stress of the steel, particularly at
temperatures below 900°C. This effect can be attributed
500 600 700 600 900 1000 1100
Temperature CC)
Fig. 5. CCC tlow stress versus temperature curves for
the materials under investigation.
678 – 40TH MWSP CONF. PROC,, 1SS 1998
ong time required to produce a significant amount of
errite for steel B, tests involving isothermal ferrite
ransformation were not performed for this steel
omposition. It was necessary to isothermally hold steels
and C for 4 and 50 minutes, respectively. This way a
icrostructure containing roughly 25 O/OVOl.ferrite could
e obtained for each.
50 , .-A.5teelA-,
o
]i lSteel B I 1
0.1 1.0 10,0 100.
Time (rein)
ig. 6. The effect of Mo and Si on the isothermal
ransformation rate of ferrite at 650”C.
3.3 Effect of bainite hold temperature on the
microstructure and retained austenite characteristics
Steels A and C were subjected to the TMP
schedule in Figure 3 to simulate different conditions for
bainite formation. Figure 7 shows the microstructure
obtained for these steels following controlled TMP and
isothermal bainite formation for 5 minutes at 300°C,
400°C and 500”C. The steels were etched using a 3°A
nital solution, revealing the sample microstructnres. The
light gray matrix phase is ferrite, the plate-like phase is a
multiphase structure of bainitic ferrite, RA and
martensite, and the black block like phases are possibly
a very fme pearlite. Careful observation and comparison
of these microstructure reveals their dependence on the
bainite hold temperature. Characteristic to both steels, the
bainite transforms from a coarse to a fine lath structure
with decreasing isothermal hold temperature. Though the
RA and martensite phases are not readily visible by the
etching technique employed, Hanzaki et al.(2’)has shown
color etching techniques to be quite effective in
distinguishing the RA from the other phases. His work
2. (a) 2. (b) 2. (c)
Fig. 7. Microstructure ofi (1) steel A and (2) steel C at various isothermal bainite hold temperatures: (a) 300°C,
(b) 400”C, (c) 500”C. (F=Ferrite, CB=Coarse Bainite, FB=Fine Bainite)
40TH MWSP CONF. PROC., 1SS 1998 – 679
revealed that the RA morphology progressively varied
from an enclosed structure surrounded by thick ferritic
bainite platelets at the higher transformation temperature
(i.e. 500”C),to a structure entrapped between laths of
bainitic ferrite (i.e. 300”C).
The Vu for each combination of steel chemistry
and TMP history (i.e. bainite treatment) was measured
using neutron diffraction. The lattice parameters and Vw
were evaluated using the Rietveld profile-refinement
method. Though the microstructure is composed of RA,
ferrite and martensite, the Vw is obtained from a
diffi-action pattern fit, which measures the relative
integrated intensities of an assumed dual phase structure
(i.e. ferrite and austenite). The slight differences in unit
cell morphology and volume ii-action between martensite
ferrite generates an overlapping and, generally,
undistinguishable and additive diffi-action pattern.
Because interest is placed on Vw, it can be assumed that
martensite and ferrite peaks are ffom a single (i.e. ferrite)
phase. Therefore, a dual phase fit of austenite and ferrite
can be confidently accomplished with little error to Vw.
Figure 8 illustrates an example of the Rietveld fit to an
experimental diffraction pattern.
Results show that the Vw is optimized at the
intermediate bainite temperature of 400”C (i.e. Figure 9),
I I 1 I 1 1 [ 1
m.
Z*
2 0
w
: -1I I
c1 I I 1 I I I 1 1 1 I
0.4 0,5 0.6 0.7 0.0 0.9 1.0 1.1
2--l>heta, deg XIOE 2
Fig. 8. Fitted neutron diffraction pattern of an
austenite/ferrite dual phase structure. (higher 2-theta
markers = austenite, lower 2-theta markers = ferrite)
hence a progressive and ongoing TRIP effect. Decreasing
the isothermal bainite hold temperature from 500 to
400”C increased the Vw in both steels, less so for steel C.
This may be attributed to the solute effect of Mo,
whereby the upper to lower bainite transformation point
is shifted to a lower temperature. Decreasing the bainite
hold temperature even further (to 300”C), on the other
hand, reversed this trend, generating lower RA levels.
where a combination of both types of RA allows for a
wide stability range against mechanical deformation and
g Steel A q steel c
t
4.7
300 400 500
Temperature ~C)
Fig. 9. Effect of bainite hold temperature on the
volume fraction of retained austenite. (time=5 min.)
680 – 40TH MWSP CONF. PROC., 1SS 1998
This can be attributed to the ease with which carbides
form at this temperature. Carbon precipitates as thin films
1.50
1.25
c 1.00
~
$
~ 0.75
0.50
0.25
n Steel A q steel c
300 400
Temperature (“C)
Fig. 10, Effect of bainite hold temperature on the C
content of retained austenite. (time=5 min.)
of carbide within the bainitic ferrite platelets, thereby
decreasing its availability for the stabilization and
retention of RA.(21)
Both steels exhibit a very similar RA carbon
(Cw) content trend with varying bainite hold temperature
(i.e. Figure 10). A maximum is observed at 400”C in both
cases. However, Hanzaki et al.(2]) and Di Chiro et al.(4)
observed an optimum C~A at 300”C for steel A with a
slightly higher Nb level (i.e. 0.035 ‘Yowt.). This
discrepancy can be explained by the different hold times
utilized (i.e. 5 min. versus 2 min.). It can be postulated
that precipitation is initiated following a critical hold
time, which when suqm.ssed, is responsible for driving the
CM as well as the Vw to lower levels. Therefore,
maximum CM occurs before 5 minutes at 300”C.
At 500”C, the CM of steel C is significantly
below that of steel A. The coarse nature of the bainite in
this steel may be impeding the stabilization of austenite
by generating a coarse RA particle and a longer range
difiision of carbon upon bainitic transformation, both of
which rendering austenite saturation with C more
difficult. This would suggest that the Vw be driven down.
Furthermore, the encapsulated RA morphology typical of
a lower bainite is inherently more stable against the SIT
by the greater load bearing nature of the surrounding,
3.4 Variation of Mechanical properties with Isothermal
Bainite Hold Temperature
The mechanical properties of the steels, which
have undergone isothermal bainite treatment, are plotted
in Figure 11. Both steels exhibit an increase in the UTS
with decreasing hold temperature, which is largely in part
due to a the finer bainitic structure,t4’21J The UTS
measurements for steel A are lower than those observed
by Hanzaki et al.(2]) and Di Chiro et al..(4) The slightly
lower Nb level maybe responsible for these lower tensile
properties. Steel C, on the other hand, generated values
within the range obtained by the previous authors. These
higher UTS measurements, compared to steel A, maybe
attributed to the strong solid solution strengthening effect
of Mo as well aa to the slightly greater Nb level.
Steel A exhibits an optimum ductility at 400”C,
which coincides with the maximum Vw. The elevated
VU, in conjunction with the wide distribution of stability
levels obtainable at this temperature, as explained earlier,
are responsible for the large ductility measurements
observed. Conversely, steel C exhibits an increasing
elongation with decreasing bainite hold temperature. This
trend, however, does not coincide with those observed for
Vw nor CM. Mo may be taking part in solid solution
(
Fig.11. Mechanical properties
isothermal bainitic treatments.
40TH
following various
time=5 min.)
MWSP CONF. PROC., 1SS 1998 – 681
high strength, fine bainite.
n Steel A q steel c
—
strengthening the RA against premature SIT, thereby
postponing fracture to larger strains.
n Steel A
+-
35.9
q steel c
T
300 400
Temperature ~C) Temperature (“C)
500
2
g Steel A
+
8.2
q steel c
[
8.8
5 10
Time (min.)
Fig. 12. Effect of bainite hold time on the volume
fraction of retained austenite. (temperature =400°C)
1.50
1.25
+ 1.00
~
$
j 0.75
0.50
0.25
2
g Steel A q Steel C
5
--E-
1.04
10
Time (min.)
Fig. 13. Effect of bainite hold time on the C content of
retained austenite. (temperature =400°C)
RA characteristics (i.e. work hardening, C content,
substitutional content).
3.6 Variation of mechanical properties with isothermal
bainite hold time
3.5 Effect of bainite hold time on the retained austenite
characteristics
Experiments veri@ng the effect of hold time on
the RA characteristics were performed at 400”C. Figures
12 and 13 illustrates the variation of Vw and Cw with
hold time, respectively. The Vw and CM for both steels
are quite similar, thereby confining the strong
dependence of the RA on the microstructure, as
processing was controlled to obtain similar structures.
Steel A exhibits a maximum Vw at 5 min.. The initial rise
of Vw with time is caused by the stabilization of
austenite through its saturation with C rejected from
adjacent forming bainite laths. Subsequent
supersaturation of the austenite phase leads to the
formation of carbides and a drop in the Vw. The Vw for
steel C, on the other hand, increases slightly with time.
Though this variation appears insignificant, the Vw of
this steel is expected to eventually drop, as the hold time
increases, for the reason mentioned above. The increased
volubility of C in austenite, as well as the strong drag
effect, in the presence of Mo, may be responsible for
postponing carbide precipitation and the drop in V~A.The
CM follows quite closely the variation in Vw for both
steels. For a given type of microstructure this is to be
expected as, the RA morphology effect remains constan~
This way stabilization becomes solely dependent on the
682 – 40TH MWSP CONF. PROC., 1SS 1998
Testing revealed that isothermal hold time has
an important effect on the mechanical properties of both
steels (Figure 14). In both instances, the UTS slightly
increased with hold time. This can be attributed to the
progressive increase in bainite being formed, as well as to
the possible precipitation of carbides. (z’2]) Ductility, on
the other hand, follows VM.(2>21)Furthermore, the
increased dispersion and refinement of austenite during
bainite transformation as time is elapsed, may also be
responsible for improving the elongation to failure by
generating a more homogeneous RA bearing structure.
This would allow for a more progressive and
homogeneous transformation.
2
--I-
1175
n Steel A
5 10
Time (min.)
2
Fig. 14. Mechanical properties followiug various
isothermal bainitic treatments. (temperatu;e=400°C)
q Steel C
.
II
T
3s.5
5 10
Time (min.)
ACKNOWLEDGEMENT
tee
atu
an
.
.
.
.
.
.
.
40TH MWSP CONF. PROC., 1SS 1998 – 683
4. CONCLUSIONS
Mo was added to Si-Mn TRIP steels to veri@
whether it can partially replace Si. This addition showed
to generate an excellent combination of strength and
ductility comparable to steels without Mo, but with
greater Si levels. The mechanical properties of these
steels were discussed in terms of their RA characteristics
and their microstructure. The results observed are
summarized:
1. CCC testing showed Mo to have a strong drag effect.
The Arl in the presence of Mo was decreased by
-35°C. Furthermore, Mo had an important solid
solution strengthening effect which was evidenced
by an accelerated increasing flow stress as the steel
was cooled.
2. Isothermal holding at 650°C prior to bainite
treatment revealed that Mo has a retardation effect
on both the formation of ferrite and cementite.
3. The mechanical properties of these steels are
dependent on the processing conditions in the bainite
region (i.e. temperature, time). The UTS is sensitive
to microstructure, whereas the total elongation is
dependent on the microstructure, and the RA
characteristics, which, also, are sensitive to the
evolving microstructure.
S
N
C
1
2
3
4
5
6
7
The authors would like to thank the Canadian
l Industry Research Association (CSIRA) and the
ral Sciences and Engineering Research Council of
ada (NSERC) for their financial support.
REFERENCES:
V.F. Zackay, E.R. Parker, D. Fahr, and R Bush,
Trans. Am. Sot. Mat., Vol. 60, 1967, pp.252.
A. Zarei-HanzaJci, Ph.D. Thesis, McGill University,
1994,
A. Zarei-Hanzaki. P.D. Hodeson and S.Yue. ISIJ
International, VOL 35, 1995,fio. 3, pp. 324. ‘
DiChiro, J. Root and S.Yue, 37ti MWSP Conf. Proc.,
~, Vol. 33, 1996, pp. 373.
W.C. Jeong, D.K. Matlock and G. Krauss, Material
Science and Engineering, A165, 1993, pp. 1.
Y. Sakum4 O. Matsumura and O. Akisue, ISIJ, Vol.
31, 1991, pp.1348.
Y. Sakuma, O. Matsumura and H. Takeshi.
Metallurgical Transactions A, Vol. 22A, 1991, pp~
489.
8. Itami, M. Takahashi and K. Ushio&, High Strength
Steels for Automotive Symp osium Proceedings,
1994, pp. 245.
9. G.B. Olson and M. Cohen, Metallurgical
Transactions, Vol. 13A, 1982, pp. 1907.
10. Tsukatani, S Hashimoto and T. Inoue, ISIJ
International, Vol. 31, No. 9, 1991, pp.992.
11, R. LeHouilier, G. Begin and A. Dobe, Metallurgical
Transactions A, Vol. 2A, 1971, pp,2645,
12. V.M. Pivovarov, LA. Tanaka, and A.A. Levchenko,
Phys. Met. Metallogr., Vol. 33, 1972, pp.1 16.
13. H. K.D.H. Bhadeshia and D.V. Edmonds,
Metallurgical Transactions A, Vol. lOa, 1979, pp
895.
14. M.G, Akben, B. Bacroix and J.J. Jonas, Acts Metall.,
Vol. 31 (1983), p. 161.
15. J.J. Jonas, Proceedings of an International
Conference sponsored by the Ferrous Metallurgy
Committee of the Metallurgical Society of the
Americam Institute of Mining, Metallurgical and
Petroleum Engineers (A. I.M.E.) and the Australian
Institute of Metals, held at the University of
Wollongong, Wollongong Australia (Aug. 20-24,
1984), p.80.
16. H. Wada and R.D. Pehlke, Metall. Trans., Vol. 16B
(1985 ),p.815.
17. A. Zarei-Hanzaki, P.D. Hodgson and S. Yue, 33rd
MWSP Conf. Proc., ISS-Aime, Vol. 29 (1992),
p.460.
18. I.P, Swainson, N.B. Konyer and J.H. Root, “1998 C2
DUALSPEC Powder Diffractometer User’s Guide”,
(1997).
19. G.E. Lucas, G,R. Odette and J.W. Sheckard, ASTM
STP 888, p.1 12.
20. A. Zarei-Hanzaki, R.Pandi, P.D. Hodgson, and
S.Yue, Metallurgical Transactions A, Vol. 24A
(1993), p. 2657.
21. A, Zarei-Hrmzaki, P.D. Hodgson, and S.Yue, ISIJ
International, VO1.35, 1995, 79-85.
684- 40TH MWSP CONF. PROC., 1SS 1998
	MAIN MENU
	PREVIOUS MENU
	--------------------------------------
	Start of Paper
	Print
	Exit CD-ROM