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ray-iim-2002-HSLA80.pdf
Archived in [email protected]
http://dspace.nitrkl.ac.in/dspace
Trans Indian Inst Metals
Vol.55, Nos. 1-2, February-April 2002, pp. 15-24
(TP 1825)
Investigations on the Microstructure and Quantification
of Mechanical Properties of a Heat Treated
Cu-bearing HSLA-80 Steel
P. K. Ray 1, R. I. Ganguly 2, A. K. Panda 2
1
- Dept. of Applied Mechanics & Hydraulics,
Regional Engineering College, Rourkela 769008, INDIA.
2
- Dept. of Metallurgical Engineering,
Regional Engineering College, Rourkela 769008, INDIA.
1
Abstract :
An ULCB steel (named GPT) received from US Naval Research Laboratory
has been characterised with respect to microstructure and mechanical properties. The
effect of heat treatment parameters, such as austenitisation temperature, tempering
temperature and tempering time has been studied. The microstructures obtained
through different heat treatment have been studied through optical, scanning electron
and transmission electron microscopes. The mechanical properties have been
quantified with respect to heat treatment process parameters. Regression equations
have been developed through 22 factorial design of experiments. These equations can
be effectively used for optimisation purposes. Also it is possible, through these
equations, to determine the heat treatment process variables for a desired
combination of mechanical properties.
Keywords : HSLA-80 steel, heat treatment, mechanical properties
1. Introduction
HSLA 80 steel has come out as a replacement of Quenched & Tempered
HY 80 grade steel. The improved technology has enabled the steel to become cleaner
and more weldable [1]. The alloy chemistry of HSLA 80 suggests that it can admit
alloying elements up to 4.5 wt %age [2]. The steel is designed with a view to having
tensile properties such as minimum yield strength ∼550 MPa with an elongation ≤
18% and a V notch Charpy value of 85 Joule at -81°C [3]. The alloy composition for
HSLA 80 steel is adjusted so as to combine the benefit of thermomechanical
processing (resulting in grain refinement) along with judicious adjustment of strength
through precipitation hardening with improved toughness at very low temperature [4].
The CCT curves for these steels indicate that there is a possibility of formation of
combined microstructures consisting of Acicular ferrite / Bainite / low carbon
martensite through controlled rolling and controlled cooling of the steel.
Owing to Cu content in the steel, this grade of steel responds to tempering after
quenching [4-7]. In the first stage of tempering (450 °C), coherent precipitates of Cu
clusters (presumably BCC) are observed. The matrix retains lath structure with
dislocations. The clusters are not identifiable in TEM microstructures due to their
fineness. The strength properties become the highest at this stage of tempering.
However, the low-temperature impact value drops significantly at low temperature.
During the second stage of tempering (500-600 °C), spherical ε-Cu (FCC) precipitates.
The strength value decreases owing to partial recovery and coarsening of the
precipitates with corresponding increase in impact values at low temperatures.
Between 600-650 °C (stage III) the Cu precipitates change their shape from spherical
to rod with appreciable recovery occurring in the microstructure [6]. During the fourth
stage of tempering 650-700 °C, second generation austenite forms, which can be
2
revealed in the microstructure [6]. It is interesting to note that there is enhancement of
strength and impact properties by tempering the steel in this region of time and
temperature of tempering [5,7]. Therefore, optimization of properties is needed
through control of process parameters such as temperature and time of tempering
within this zone. Though some attempts have been made to optimize the properties
through single factor experiments [5,6], less attempt has been made to quantity the
effect of the variables on the strength and toughness properties. The present study
attempts to quantify the effects of the processing parameters on mechanical properties
such as YS, UTS, and Charpy impact values at low temperature to enable selection of
the processing variables for the desired properties. Whenever necessary, structural
characterization has been done by Optical, Scanning, and Transmission Electron
Microscopy.
2. Experimental
The steel, designated GPT, was received from US Naval Research Laboratory
(see Table 1 for composition) and is designated GPT. It was characterized with
respect to inclusion content and grain size using Quantimet 570 image analysis system.
AC1 and AC3 temperatures were measured using DTA, TMA and Dilatometer (Table
2). The heat treatment parameters studied were (i) austenitisation temperature, (ii)
tempering temperature, (iii) tempering time.
Table - 1 Chemical composition of GPT steel
C
0.05
Mn
1.00
P
0.009
S
0.001
Si
0.34
Cu
1.23
Ni
1.77
Cr
0.61
Mo
0.51
Al
0.025
Cb
0.037
V
0.004
Ti
0.003
Table - 2.
Steel
GPT
Ac1 and Ac3 values (Calculated and Experimental).
Ac1 (°C)
Ac3 (°C)
Calculated Experimental Calculated
Experimental
702
690
899
860
3
Temperatures were maintained within the accuracy limit of ±3°C.
Microstructures of all the samples were examined by optical, scanning and
Transmission Electron Microscopes. Fractured samples (Tensile & Charpy) were
examined in Scanning Electron Microscope. Mechanical properties were determined
in Instron testing unit (model 1195) for the as-received as well as heat treated samples
adopting ASTM E-8-78 method.
In order to determine the region of optimum properties, heat treated samples
were subjected to Charpy impact test at -50 °C. Charpy values were higher in the
neighborhood of tempering temperature of 700 °C where the strength values were
adequate i.e., YS ≥ 600 MPa or 85 ksi. It was therefore decided to design experimental
matrices in this region and to form regression equations around tempering temperature
of 650 °C.
3. Results & Discussions
The average grain size of the as-received steel was found to be between 7 to
8µm. The inclusion content of the steel was measured to be very low (<10-4 Area
fraction). This is attributed to very low content of P and S in the steel. The
mechanical properties of the steel were found to be as follows:
YS = 636.7 MPa , UTS = 708 MPa ,
% Elongation = 30% on 25mm GL
Fig 1 shows Scanning Electron Microphotograph of as-received GPT steel. The
microstructure indicates Bainitic / non-polygonal or acicular ferritic structure. The
fractrograph of as-received tensile sample shows typical dimpled structure (Fig 2).
Fig. 1 - SEM of As-received steel
Fig. 2 - Fractograph of broken tensile
specimen of As-received steel
4
Fig 3 shows the effect of tempering temperatures on hardness values of steel
austenitised at 900, 950, 1000 °C for 1 hour followed by quenching in water and
cooling in air.
380
900 oC-AC
360
900 oC-WQ
950 oC-AC
Hardness (VPN)
340
950 oC-WQ
320
1000 oC-WQ
300
280
260
240
220
400 450 500 550 600 650 700 750
Tempering Temperature oC
Fig. 3. Hardness vs. tempering temperature curves
It may be observed from Fig. 3 that austenisation temperature did not alter the
tempering behaviour of the steel either in water quenched or in air cooled condition.
However, there is increase in the as quenched hardness values of the steel with
decreasing austenisation temperature. This is due to increased amount of retained
austenite in quenched sample owing to austenitisation of the steel at higher
temperature. Similar trend was not observed for the steel cooled in air from different
austenisation temperatures. The hardness values of the water quenched steel were
ranging between 310 - 340 VPN for samples austetinised at different temperatures,
whereas the hardness values of the air cooled sample were nearly the same, i.e., 260
VPN at all austetinisation temperatures. The difference in the hardness values of
water quenched and air cooled samples is attributed to the difference in the
microstructures obtained as a result of different rates of cooling (comparison between
Figs. - 4 and 5). As seen from Fig. 4, the microstructure for water quenched steel is
acicular in nature, whereas Fig. 5 shows non-polygonal ferritic / Bainitic
microstructure in the air cooled sample.
5
Fig. 4 - SEM of WQ steel
Fig. 5 - SEM of AC steel
The hardness vs. tempering temperature curves show peak values in the
vicinity of 450 °C for water quenched samples. Similar peak is also observed for air
cooled samples with slight shift of the peak towards higher tempering temperature.
The shift of peak for air cooled sample possibly indicates retardation in the kinetics of
precipitation for air cooled structure. The probable cause for slight change in the peak
temperature is attributed to the difference in the initial microstructure of the water
quenched and air cooled samples. Since water quenched and air cooled steel show
similar tempering behaviour, it was decided to study quantitative effect of tempering
parameters on mechanical properties of water quenched steels only.
The transmission electron microphotographs of the as-quenched (water
quenched from 900 °C, 1 hour) samples (Figs. 6a,b) show matrix of lath martensite
containing dislocations. On an average the width of the laths was ∼0.4 µm as
measured in the microscope itself. The areas of austenite region were verified using
SAD (typical FCC ring pattern). The region of austenite was further confirmed from
dark field image. The continuous rings of SAD pattern were indicative of fineness of
the austenite grains. These austenites are the retained austenite in the quenched sample
as was observed by various workers [6].
Fig. 6a - TEM of WQ steel showing
matrix of lath martensite
Fig. 6b - TEM of WQ steel
showing dislocations
6
Fig. 7a - TEM of Quenched and tempered
steel (450 °C, 1 hour) showing
lath structure with dislocations
Fig. 7b - TEM of Quenched and
tempered steel (450 °C,
1 hour)showing
Figs. 7a,b show the transmission electron microphotographs of quenched and
tempered (450 °C, 1 hour) samples. Fig. 7a shows lath structure with dislocations. Fig.
7b is the TEM studies of the same at some other region showing fine precipitates
occurring due to tempering.
Fig. 8 - TEM of Quenched and tempered steel (600 °C, 1 hour)
showing recovered structure
Tempering at 600°C causes recovered structure. TEM photograph (Fig. 8)
shows there is less dislocation in the microstructure in white ferritic matrix.
Tempering at 700°C resulted recovery of the dislocations in the laths with some
precipitates (Fig. 9a). Dark region of second-generation austenite is seen at the lath
boundaries (Fig. 9b). The darkness of the austenite is due to quenching of the steel
after tempering at 700 °C [5,6].
7
Fig. 9a - TEM of Quenched and tempered
steel (700 °C, 1 hour) showing recovery
of dislocations with precipitates
Fig. 9b - TEM of Quenched and
tempered steel (700 °C, 1 hour)
showing second generation austenite
(dark regions) at lath boundaries
Since the steel is designed primarily as a structural material for ship building
purpose, the requirement for low-temperature impact properties is very stringent.
Therefore, it is necessary to evaluate Charpy values of different heat-treated steels at
low temperature before optimization of properties is taken up. Charpy samples used
for the above purpose were heat treated as described in Table 3.
Table - 3 . IMPACT VALUES at -50 °C
Treatment
-50 °C CVN value (J)
As recieved
244
122
(950°C 1hr. - WQ)
Air Cool
149
102
WQ + tempered at 450°C for
1hr.
203
WQ + tempered at 600°C for
1hr.
258
WQ + tempered at 700°C for
1hr.
While the Charpy value was the least (102 Joules) for the sample tempered at
450 C for 1 hr. it was maximum (258 Joules) for the sample tempered at 7000C for 1
hr. Figs. 10 and 11 represent fractured surfaces of impact test samples tempered at 450
°C and 700 °C respectively. It is evident that while the fractrograph of the former
sample (tempered at 450 °C) shows river pattern with quasicleavage areas (typical
brittle fracture) (Fig. 10), the later sample (tempered at 700 °C ) shows dimples,
indicative of typical ductile fracture (Fig. 11 ).
0
8
Fig. 10 - Fractograph of broken Charpy
specimen (Quenched and tempered at
450 °C, 1 hour) shows quasicleavage nature
Fig. 11 - Fractograph of broken Charpy
specimen (Quenched and tempered at
700 °C, 1 hour) shows dimple fracture
Analysing the results of the impact properties obtained for the heat treated
steel, extensive experiments were planned to evaluate mechanical properties of the
steel between tempering temperature of 600-700 °C for wide range of tempering time.
It may be mentioned here that beyond the tempering temperature of 650 °C, hardness
increases (Fig.3). The increase in hardness beyond 650 °C may be explained by the
fact that the Ac1 temperature for this steel being less than 700 °C (Table - 3), the
austenite formed at the higher temperature range of tempering, changed to martensite
on quenching after tempering (vide dark regions at the lath boundaries in Fig.9b).
Further the increase in toughness in this tempering temperature range is attributed to
the formation of new generation austenite [6]. Therefore, in the region of tempering
temperature of 600 - 700 °C both the impact value and strength value reach optimum
condition best suited for use in Naval environment.
Table 4 - Tensile Properties of GPT steel after tempering.
G P T steel
Temp °C /
Time (hr)
700 / 0.33
700 / 2
700 / 12
700 / 80
600 / 0.33
600 / 2
600 / 12
AQ
YS
MPa
556.2
483.6
413.0
232.5
843.7
708.3
661.2
867.2
TS
MPa
662.2
633.7
570.9
341.1
853.5
715.1
673.9
954.5
YS/TS
N
0.85
0.76
0.72
0.68
0.99
0.99
0.98
0.91
0.19
0.21
0.25
0.19
0.078
0.10
0.09
-----
9
Table 4 shows the tensile properties of the steel in different heat-treated
conditions. It is seen from the Table that Y.S/T.S ratio of the steel tempered at 7000C
reach a value between 0.68-0.76 whereas for the steel tempered at 600 °C this value
remains constant, i.e., 0.98-0.99. Fig. 12 compares typical stress-strain diagram of the
steel tempered at 600 °C and 700 °C respectively for 1hr. The stress - strain diagram
of steel tempered at 600 °C, shows discontinuous yielding. In contrast, the stressstrain diagram of the steel tempered at 700 °C shows continuous yielding behaviour,
typical of a Dual phase steel.
Fig. 12 - Stress-strain diagram of specimens quenched and
tempered at 600 °C and 700 °C
10
900
YS 700 oC
Yield Stress (MPa)
800
YS 680 oC
700
YS 650 oC
600
YS 620 oC
500
YS 600 oC
400
300
200
15
16
17
18
19
Holloman-Jeffe (K-hr x10^3)
20
Tensile Stress (MPa)
Fig. 13 - YS vs. Hollomon-Jaffe Parameter
900
TS 700 oC
800
TS 680 oC
700
TS 650 oC
600
TS 620 oC
500
TS 600 oC
400
300
200
15
16
17
18
19
20
Holloman-Jaffe (K-hr x10^3)
Fig. 14 - UTS vs. Hollomon-Jaffe Parameter
Figs. 13 and 14 show the plot of Holloman-Jaffe temperature normalised
parameter with YS and UTS respectively. The YS values fitted well in a straight line
with the Holloman-Jaffe parameter at all tempering temperatures as was observed by
others [5]. However there is a marked deviation from straight line fit for UTS values
while plotted against the parameter, particularly beyond 650 °C of tempering
temperature. The possible explanation is attributed to the formation of austenite during
tempering above 650 °C. This austenite forms martensite while quenched from
tempering temperature and thus forming dual phase microstructure.
The YS and UTS quantified with respect to time and temperature of tempering
by Best-fit method are as follows :
YS = ( 2305.105 - 2.5604 T ). t(3.4589 - 0.551 ln T)
--------- ( 1 )
TS = ( 1528.94 - 1.341 T ). t(1.994 - 0.322 ln T)
11
----------- ( 2 )
The equations can be utilised for finding the mechanical properties of the steel
for different combinations of time and temperature of tempering. Similar equation is
obtained for Charpy value of the steel at -50 °C and is shown below in equation (3) :
CVN = (0.03544 T - 24.3566).t2 + (-0.52484 T + 357.7386).t + (1.0706 T -493.13)
................ ( 3 )
In eqns. (1) to (3), T is tempering temperature in °C and t is the time of tempering in
hours. YS and TS are the yield strength and tensile strength in MPa, CVN is the -50 °C
Charpy impact value in Joules.
Since it is observed from Tables 3 and 4 that the best combination of properties
(YS and -50 °C CVN) occurs in the vicinity of 700 °C and 1 hour time, it was
therefore decided to carry out 22 factorial design of experiments in this region of time
and temperature of tempering. Table 5 shows design matrices for YS and -50 °C CVN
value prepared over a tempering temperature range of 600-700 °C and tempering time
of 0.33-2 hours. Regression equations for YS and CVN value, which are formed from
the design matrices, give quantitative estimate of the properties and can be effectively
used for optimization purposes [9,10].
YS = 647.95 – 128.05 X1 - 52 X2 + 15.7 X1. X2
CVN = 213.75 + 26.75 X1 + 15.25 X2 - 21.75 X1. X2
......................... ( 4 )
..........................( 5 )
..........................( 6 )
where X1 = (x1 - 650) / 50, X2 = (x2 - 1.165) / 0.835
x1, x2 are natural values of temperature and time of ageing and YS is the yeild
strength (MPa) and CVN is the Charpy V-notch value at -50 °C (Joule). X1 and X2
are in coded form and can be decoded by using the relations given in eqn. (6).
Table 5 - 22 Design Matrix along with responses
INPUT VARIABLES
Ageing
Ageing
temperature °C
(hrs)
RESPONSES
time
YS (MPa)
-50 °C
(Joule)
600
0.33
843.7
150
600
2.00
708.3
224
700
0.33
556.2
247
700
2.00
483.6
234
12
CVN
The -ive coefficients of X1 in eqn.(4) implies that the yield strength will
decrease for an increase in tempering temperature above the base level (650 °C ⇒
+ive X1). The +ive coefficient of X2 in eqn.(5) indicates that there will be increase in
Charpy value for +ive X1. However if X1⇒0+, then the decrease in YS will be
negligible. For decrease in tempering time below the base level (1.165 hours ⇒ -ive
X2), the yield strength increases since the coefficient of X2 in eqn.(4) is –ive. The
Charpy value decreases since the coefficient of X2 in eqn.(5) is +ive. However this
decrease is not much since the numerical value of the coefficient of X2 in eqn.(5) is
small and X2 is a fractional quantity. The effect of the time-temperature interaction is
small in both cases since the coefficients of (X1.X2) are small in eqns. (4) and (5), and
two fractional quantities X1 and X2 are multiplied. Thus by tempering the steel at a
temperature just above the base level (650 °C, i.e., X1⇒0+) and time near the end of
the range (⇒0.33 hours, i.e., X2⇒1-), there will be some increase in yield strength
with high increase in toughness. Thus there is distinct advantage of tempering at a
temperature above the base level temperature and time below the base level time.
4. Conclusions
1) The steel responds to tempering both in air cooled and water quenched conditions.
2) Austenitisation at 900 °C followed by quenching in water produces martensitic
structure which is distinctly different from the microstructure of austenitised and
air cooled steel.
3) Presence of retained austenite is confirmed by SAD and dark field image of TEM
microstructure.
4) Maximum strength value and minimum Charpy value at -50 °C are observed by
tempering the steel at 450 °C after quenching or air cooling the steel from the
austenitisation temperature of 900 °C.
5) The steel shows high value of YS/UTS ratio and discontinuous yielding behaviour
after tempering at 600 °C. However, tempering at 700 °C results decreasing
YS/UTS ratio with continuous yielding behaviour, typical of Dual Phase steel.
6) The YS and UTS values were plotted against Holloman-Jaffe temperature
normalised parameters. The plots of Yield stress vs. Hollomon-Jaffe Parameter
show linearity within wider range of tempering temperature. However, similar
plots for UTS vs. Hollomon-Jaffe Parameter show significant deviations from
linearity at higher tempering temperature. This is due to the formation of austenite
at higher tempering temperature.
7) YS and Charpy values have been quantified with respect to time and temperature of
tempering in the range of 0.33-12 hours and 600-700 °C respectively by using Best
fit method.
8) The quantitative effect of tempering parameters (i.e. temperature and time of
tempering) varying in a shorter range are shown in the form of regression
13
equations obtained by applying statistical Design of Experiments. These
quantitative relations are utilised for optimization purpose.
5. References
1. Montemarano T.W., Sach B.P., Gudas J.P., Vassilaros M.G. and Vandervelt H.H.,
High Strength Low Alloy Steels in Naval Construction, J. Ship Production, 3:145162, 1986.
2. Steel Plate, Sheet, or Coil, Age-Hardening Alloy, Structural, High Yield Strength
(HSLA-80 and HSLA-100), January 1990, MIL-S-24645A(SH).
3. Coldren A.P. and Cox T.B., Development of 100 ksi Yield Steel, Technical Report,
David Taylor Research Laboratory, 1985, DTNSRDCN00167-85-C-006.
4. Wilson A.D., Hamburg E.G., Colvin D.J., Thompson S.W. and Krauss G.,
Properties and Microstructures of Copper Precipitation Aged Plate Steel,
Microalloying ‘88, Chicago, ASM, September 1988.
5. Foley R.P. and Fine M.E., Microstructure and Property Investigation of Quenched
and Tempered HSLA-100 Steel, Proc. of Int. Conf. on Processing, Microstructure and
Properties of Microalloyed and Other Modern High Strength Low Alloy Steels,
Pittsburgh PA, Iron and Steel Society, June 3-6, 1991.
6. Mujahid M., Lis A.K., Garcia C.I. and DeArdo A.J., Structure-Property Studies of
Cu-containing HSLA-100 Steels, Proc. of Int. Conf. on Processing, Microstructure and
Properties of Microalloyed and Other Modern High Strength Low Alloy Steels,
Pittsburgh PA, Iron and Steel Society, June 3-6, 1991.
7. Goodman S.R., Brenner S.S. and Low J.R. Jr., An FIM-Atom Probe Study of the
Precipitation of Copper from Iron-1.4 At. Pct. Copper, Part I: Field Ion Microscopy,
Metallurgical Transactions, 4A:2363-2369, 1973.
8. Mikalac S.J. and Vassilaros M.G., Strength and Toughness Response to Aging in a
High Copper HSLA-100 Steel, Proc. of Int. Conf. on Processing, Microstructure and
Properties of Microalloyed and Other Modern High Strength Low Alloy Steels,
Pittsburgh PA, Iron and Steel Society, June 3-6, 1991.
9. Panda A.K., Ganguly R.I. and Misra S., Tool and Alloy Steels, March/April
(1979), pp 101-108.
10. Hicks C.R., Fundamental Concepts in the Design of Experiments, Holt Rinehart
and Winston Inc., New York, (1964).
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