1. Experience in Producing
Vanadium-Microalloyed Steels
By Thin-Slab-Casting
Steel Technology
By
Robert J. Glodowski
Director of Technical Services
Stratcor, Inc.
A Subsidiary of Strategic Minerals Corporation
Pittsburgh, Pennsylvania; U.S.A.
Presented at:
TSCR 2002
The International Symposium on Thin-Slab Casting and Rolling
Guangzhou, China
December 3-5, 2002

2. Experience in Producing Vanadium Microalloyed Steels
By Thin-Slab-Casting Steel Technology
By Robert J. Glodowski
ABSTRACT Because many of these new mills did not have
Vanadium-nitrogen microalloying has proven to be downstream finishing processes such as cold rolling
highly compatible with the thin-slab casting and the and coating, there was a natural commercial interest in
direct-charging steelmaking process. Metallurgical fea- developing applications for as-rolled hot band. Since
tures of vanadium and nitrogen are related to the unique much of the market for hot band was for structural
process parameters associated with this process. and automotive applications, the demand for higher
Key words: Thin slab, vanadium, nitrogen, solu- strength grades was a natural extension of this hot-
bility, precipitation, grain size, strain aging. band market. Conventional microalloying approaches
for high-strength, low-alloy (HSLA) strip steels pre-
INTRODUCTION: GROWTH OF THIN-SLAB sented some processing complications that were new
CASTING AND DIRECT CHARGING and required some alternative solutions2. For many
The thin-slab-casting revolution has taken place mills, the vanadium and nitrogen alloying system has
over a relatively short time period, starting first in proven to be highly effective and compatible with the
Crawfordsville, Indiana, U.S.A. in 1989. The immedi- processing peculiarities of the thin-slab-casting and di-
ate success of that operation led to many more plants rect-charging process.
being built, first in North America and then spreading
across the world1. Initially, the emphasis was on the PROCESSING CHARACTERISTICS:
obvious economic and ecological benefits gained by VARIATIONS FROM NORMAL PRACTICES
retaining part of the heat of casting in the slabs through The thin-slab-casting, direct-charging and finish-roll-
to the rolling process, and casting to a “near net shape,” ing process has a number of characteristics that vary
reducing the amount of hot working necessary to re- from conventional blast-furnace / BOF / slab-cast /
duce the material to final gauge. Using electric-furnace slab-reheat / roughing-mill / finishing-mill routings. First,
steel with most of the iron units derived from scrap, steelmaking often uses electric-arc-furnace (EAF) melt-
the capital costs for building these types of mills was ing from predominately scrap charges, and the refin-
less than half that of conventional mills using blast fur- ing process takes place in separate ladle-furnace op-
naces and basic-oxygen furnaces. erations. While this is not true in all cases, it is by far
Because of the rapid industrial expansion of this the majority, especially for greenfield sites. Second,
new strip-manufacturing process, the effect of the sig- the casting process involves rapid cooling, necessary
nificant metallurgical changes of thin-slab casting was for high-volume production in thin slabs from 50 to
initially managed on a real-time basis by those metal- 100 mm in thickness. Third, the slab is directly charged
lurgists left with the responsibility of providing a sale- into a tunnel holding furnace without undergoing a aus-
able quality product off the mill. Laboratory-based met- tenite (γ) to ferrite (α) transformation. Fourth, the tunnel
allurgical research projects quickly followed, with a large furnace is limited in reheating capabilities, typically from
number of industrial and academic organizations evalu- to 1100 to 1150°C maximum.
ating the problems and promises of the newly devel- Finally, the 50-mm slab will enter directly into a
oped and economically successful thin-slab casting and 5- or 6-stand finishing mill without undergoing a rough-
direct-charging process. A basic understanding of the ing mill reduction. Also, the 80-to 100-mm slab cast-
impact of these processing differences was rapidly de- ers will have one or two roughing mills prior to a 5-
veloped through these projects. stand finishing mill. In those cases, a transfer system
of coil boxes or holding furnaces will be installed after
Robert J. Glodowski is Director of Technical Services; the roughing mills to equalize temperatures before en-
Stratcor, Inc., a Subsidiary of Strategic Minerals Corpora- tering the finishing-mill stands. This step provides a
tion; Pittsburgh, Pa.; U.S.A. speed “break” between the start and finish rolling
2

3. speeds to allow more reasonable entry speeds in the The carbon pickup from the ferromanganese, as
first reduction pass. The remainder of the process, well as other alloys, must be accounted for at the time
from finish rolling through a runout table usually of tapping. Melt-in carbon levels are kept low to mini-
equipped with significant accelerated-cooling capabil- mize time required to finish the decarburization pro-
ity and into a down coiling system, will be similar to cess. Nitrogen can be removed from the steel bath by
existing modern strip mills. Many of the mills have the absorption in the CO bubbles formed during the de-
capability to approach 1-mm final thickness, although carburization process. Because the amount of decar-
most commonly the final gauge will be 1.5-mm or burization is limited in a low-carbon bath, the nitrogen
larger. introduced from the scrap charge and the electric-arc
is not easily removed. Therefore, nitrogen levels re-
MELTING CONSIDERATIONS: main relatively high.
CONTROLLING RESIDUAL ELEMENTS While blast-furnace and basic-oxygen-furnace
With EAF scrap-based melting, the normal issues melting may routinely result in nitrogen levels from 30
of metallic and nitrogen residuals have to be consid- to 50 ppm, electric-arc-furnace melting may average
ered. High-quality scrap, along with pure iron units from 70 to 100 ppm. Even though some mills have
like DRI and HBI, are used for melting to minimize developed the capability to produce EAF steels at sub-
metallic residuals. Early surface-quality problems from stantially lower nitrogen levels using selected raw ma-
hot shortness were identified in some mills. Maximum terials and special processing, nitrogen levels continue
copper levels for these steels were determined for each to be a major difference between conventional BOF-
mill according to their process capabilities. Copper is melted and EAF-melted thin-slab strip steels. Even with
generally restricted to less than 0.15% in these mills the use of aluminum, considerable free nitrogen will
today. remain in the hot-band steel.
Aluminum levels are controlled to levels less than Additions of the microalloying elements -- vana-
0.035%1, 3. Higher levels are not needed for deoxida- dium, niobium and titanium -- all eventually combine
tion and can help form inclusions which can interfere (precipitate) with the interstitial elements, carbon and
with the fluid flow in the caster. Calcium treatment is nitrogen. The effect of these microalloys on the steel
also commonly used to reduce the clogging tendency properties is to a great extent determined by the type
of aluminum inclusions and to modify oxides and sul- and timing of that precipitation. For this reason, when
fides so that they remain globular through the rolling microalloys are added to steel, nitrogen becomes an
process. These modifications improve the isotropic integral part of the alloy system as much as carbon.
nature of the product, particularly the formability and Significant differences in residual nitrogen will play
toughness in the transverse direction. a large role in defining the performance of a given
Sulfur control has proven to be critical for thin- microalloy addition. Managing the nitrogen level there-
slab casting. Sulfur control can be a problem for either fore becomes a necessary part of producing consistent
basic-oxygen or electric-arc furnace melting, but well- microalloyed-steel properties. Because of the higher
established techniques are available to reduce the sul- nitrogen levels inherent in electric-arc-furnace steels,
fur levels. With the EAF process, this is done both in the microalloy used should be compatible with the ad-
the furnace and later in the ladle using calcium treat- ditional nitrogen.
ments described above. Each mill will have estab-
lished specific control levels they believe necessary CASTING CONSIDERATIONS:
for casting a quality product. Maximum sulfur levels RAPID SOLIDIFICATION
are generally below 0.01% 3. From a physical-metallurgy standpoint, the defin-
Nitrogen levels are characteristically quite high from ing characteristic of the thin-slab-casting process is the
EAF’s when producing a low-carbon melt (below rapid solidification rate necessary to successfully cast
peritectic) preferred for the thin-slab casters. While this product. In conjunction with the steelmaking modi-
many thin-slab-casting operations will claim to have fications discussed previously, rapid solidification re-
the capability to produce steels in the peritectic region sults in fine, evenly dispersed globular inclusions. While
(nominally from 0.08 to 0.16% C), most would prefer some refinement of the as-cast austenitic structure was
not to be casting in this carbon range. Staying below anticipated due to rapid solidification, experience has
the peritectic while casting requires restricted tapping shown that the grain-size remains large -- on the order
carbon levels, particularly for high-strength steels where of 1 mm -- compared to 2-3 mm for conventional
high-manganese levels are required. casting. This effect is shown in Figure 1.
3

4. DISTRIBUTION OF AUSTENITIC GRAINS EVOLUTION OF AUSTENITE GRAINS
30 1
Near Surface 0.9 1s 2s
Fraction Recrystallized
25 0.8
0.7
Frequency, %
20
0.6
1s 2s
15 0.5
1s
0.4 2s
10
Inter-Pass Time
0.3
Center 0.2
5
0.1
0 0
0 500 1000 1500 2000 2500 3000 400 600 800 1000 1200 1400
Grain Size, µm Initial Austenitic Grain Size, µm
Figure 1 -- As-cast austenite-grain-size distribution in Figure 2-- Austenite evolution during hot rolling of 50-
an industrial thin-slab sample10. mm C-Mn steel4.
REHEATING AND ROLLING with an as-cast microstructure, hot ductility of the slab
CONSIDERATIONS: GRAIN REFINEMENT becomes an issue. Ductility losses from grain-bound-
AND RECRYSTALLIZATION ary precipitates or grain-boundary penetration of liq-
Practical design considerations limit the reheating uid Fe-Cu alloy can promote hot cracking during the
capability of tunnel furnaces used to transfer the slab initial rolling pass of the dendritic as-cast structure.
from the caster to the rolling mill. Operational consid- These grain-boundary ductility issues are aggravated
erations also dictate that the residence time in the fur- by the large austenite grain size along with natural den-
nace is restricted and subject to variation. Hence, the dritic segregation during solidification. While higher ini-
more appropriate name for the tunnel furnace in a thin- tial rolling temperatures would be beneficial in mini-
slab process is a holding or transfer furnace, rather mizing hot-ductility problems during rolling, the lim-
than a reheating furnace. The primary operational func- ited reheating capabilities of the tunnel furnaces re-
tion of the tunnel furnace is to retain the heat in the strict this option.
slab as it leaves the caster and to equalize the tem- Figure 3 shows the effect of various microalloying
elements on the ductility loss of steels reheated to near
perature at a consistent level throughout the slab. The
melting to simulate the as-cast dendritic microstruc-
maximum reheating-temperature capability for tunnel
ture5. The key to preventing transverse cracking dur-
transfer furnaces is typically between 1100 and 1150°C.
ing casting or edge cracking during rolling is to main-
Unlike conventional slab processing, the steel slab tain the temperature during deformation above the point
is not cooled through a transformation from austenite where the ductility loss begins. Figure 4 shows the
to ferrite and then reheated back to austenite in the effect of the reheating temperature on the rate of for-
reheating furnace. The benefit of austenite-grain re-
finement from the γ−α and α−γ transformations dur- EFFECT OF TEMPERATURE ON DUCTILITY
ing conventional-slab reheating is therefore not avail- 100
able. This necessitates that all austenite-grain refine- 90
C-Mn
Reduction of Area, %
ment and homogenization from the initial dendritic as- 80
cast microstructure must occur during the rolling pro- 70
V
cess. Figure 2 show the results of a model4 predicting VN
60
the austenite evolution during hot rolling of a 50-mm
50
Nb
thick slab of C-Mn steel at 1150°C, near the maxi-
mum temperature capability of most thin-slab opera- 40
tions. For a starting austenite grain size of 1000 µm, a 30
50% reduction is necessary to complete recrystalliza- 20
tion in 1 to 2 seconds. At lower temperatures, more 700 800 900 1000
time or more deformation is required to complete re- Temperature, deg. C
crystallization. Figure 3 -- Effect of temperature on hot ductility of vari-
Because the initial rolling deformation takes place ous microalloyed steels5.
4

5. EFFECT OF REHEATING TEMPERATURE • Minimal impedance to static recrystallization dur-
80
ing rolling.
Rate of Formation of Copper
Layer (cm2 / s x 108) 60
0.35% Cu • Precipitation for strengthening should occur after
0.25% Cu finish rolling.
40 • Promote fine ferrite grains without need for roll-
0.20%Cu ing below the Tnr temperature.
20
• Thermal and mechanical processing requirements
0.10%Cu should not cause undue production complications.
0
(Reduction schedule, finish-rolling temperatures,
-20 and coiling temperatures should be similar to car-
bon steel.)
-40
900 1000 1100 1200 1300
Reheating Temperature, deg. C
CHARACTERISTICS OF AVAILABLE
MICROALLOYING ELEMENTS
Figure 4 -- Effect of reheating temperature on Fe-Cu Solubility Considerations
layer formation5.
For the desired characteristics for microalloys in
mation of the Fe-Cu alloy layer, which contributes to the thin-slab process, the solubility in austenite is a
hot shortness. Unfortunately, the temperature for the critical factor. Figure 5 shows the relative solubility of
maximum rate of formation of the grain-boundary pen- the carbides and nitrides of V, Nb, Ti, and Al. For the
etrating Fe-Cu alloy is the normal operating tempera- alloys V and Ti, the carbide form is considerably more
tures of the tunnel-transfer furnaces, 1050 - 1150°C5. soluble than the nitride form. For Nb, the difference is
For optimum precipitation strengthening, the less although the nitride is still less soluble at equiva-
microalloy should also be in solution in the austenite lent concentrations of C and N. Vanadium carbide (VC)
as long as possible during the rolling process. Maxi- is dramatically more soluble than any of the other car-
mum precipitation strengthening occurs either in the bides, and vanadium nitride (VN) is more soluble than
interphase region during transformation, or even more the other nitrides.
effectively, in the ferrite after final transformation. Pre- Figure 6 on the next page shows the relative
cipitates that form in the austenite are generally larger amounts of V, Ti, and Nb combined (precipitated) as
and more dispersed, reducing the strengthening con- an M(C,N) for a range of temperatures for a typical
tribution of the precipitation. multiple microalloyed steel composition. Over the typi-
Therefore, the solubility of the microalloy is sig- cal slab temperature in a tunnel furnace (1050°C to
nificant. The microalloy must be in solution in the re- 1150°C), the Ti will be fully precipitated, the Nb pre-
heating step in order to contribute to precipitation cipitation will vary widely depending on temperature
strengthening in the final as-rolled product. It should
remain in solution throughout the rolling process for MICROALLOY PRECIPITATES
maximum precipitation strengthening and to minimize Temperature, deg. C
deformation resistance caused by precipitation in the 1,400 1,200 1,000
1
austenite that can lead to increased rolling forces, dif-
0
ficulty in controlling shape, and roll wear. VC
-1
TiC
-2
VN
PREFERRED MICROALLOY CHARACTERISTICS NbN NbC
-3
Log Ks
FOR THIN-SLAB CASTING
-4
Based on the previous discussion, the preferred
-5 Liquid AlN
characteristics of a microalloy for use in a thin-slab
-6
casting process can be summarized as follows:
• Compatibility with EAF steel, particularly with
-7
TiN
-8
higher nitrogen levels.
• Minimal precipitation during solidification for en-
-9
0.55 0.6 0.65 0.7 0.75 0.8 0.85 0.9
hanced castability. 103 / T, K-1
• Highly soluble in austenite at tunnel-furnace tem- Figure 5 -- Solubility product of microalloy carbides and
peratures. nitrides 12.
5

6. MICROALLOY PRECIPITATES SOLUTE-RETARDATION PARAMETERS
50
250
0.1%C
Solute-Retardation Parameter
0.3% Si
Ti, Nb, and V in M(C3N) and
40 200
1.4% Mn
0.01% Ti
Al in AlN, wt, %
0.03% Nb 150
30 Nb M(C3N)
0.035% Al
0.08% V 100
20 0.0105%N
50
Al Ti
10
0
V V Mo Ti Nb
0
Microalloying Element
800 1000 1200 1400
Figure 7-- Solute-retardation parameters for static re-
Temperature, deg. C
crystallization (per 0.01%)8, 9.
Figure 6 -- Pricipitation of carbonitrides in microal-
loyed steel6. peratures must offset any recrystallization retardation
caused by microalloys.
but will never be fully in solution, and the V will al- Uranga10 noted that decreasing Nb content nota-
most all be in solution. The preferred condition, as bly improved the microstructural homogeneity of thin
discussed earlier, is that the microalloy be in solution slabs by reducing the fraction of coarse grains and
in the austenite at the start of rolling. Only the vana- also the size of the largest grains. Smaller reductions
dium will meet the criteria of being in nearly full solu- in Nb could be tolerated if the entry rolling tempera-
tion for the conditions typical for rolling thin slabs. ture were increased. Ti has an intermediate SRP value
between V and Nb. However, TiN precipitates will
Recrystallization Considerations exist in the as-cast microstructure and will likely in-
Microalloying additions should not significantly im- hibit recrystallization. Vanadium, because of its high
pede the necessary static recrystallization needed after solubility and low SRP value, has little effect on the
the first reduction pass. Static recrystallization can be recrystallization process during the initial rolling phase.
inhibited by precipitates that are present in the as-cast Thus, the combination of high solubility of the V(C,N)
austenite on entering the first roll stand, or by strain- precipitates and the low contribution of V in solution
induced precipitates that form immediately from the to solute drag make vanadium the desirable micro-
first rolling process7. As discussed previously, vana- alloying element for minimal retardation of the neces-
dium has the best chance of staying in solution during sary recrystallization and homogenization of the den-
this initial rolling stage because of the higher solubility dritic as-cast austenitic microstructure.
of both the nitrides and carbides compared to the al-
ternative microalloys. Recrystallization Considerations
Solute drag from the microalloy in solution can As discussed earlier, the nitrides of V, Ti and Nb
also retard static recrystallization of the as-cast struc- tend to form before the carbide precipitates. Vanadium
ture. Figure 7 shows solute retardation parameters is unique in that nitrogen-rich V(C,N) is the preferred
(SRP) for various microalloying elements8, 9. These precipitate for effective ferrite strengthening. In the
SRP values quantify the delay produced in recrystalli- case of Nb, nitrogen will encourage the strain-induced
zation time by the addition of 0.01% of each alloy precipitation of nitrogen-rich Nb(C,N) early in the roll-
element in solution with respect to a C-Mn base steel. ing process. Nb is then removed from solution that
As shown, vanadium has the least effect of retarding would otherwise be available for precipitation strength-
recrystallization, and niobium has the greatest effect. ening of the ferrite8. The maximum effective Nb addi-
The potential for a microalloying element to delay re- tion level is limited by the maximum reheat tempera-
crystallization is critical because of the necessity of ture available, nitrogen levels, and the ability to achieve
insuring static recrystallization after the first reduction sufficient recrystallization of the dendritic as-cast struc-
pass. Either additional reduction and/or higher tem- ture. Hensger and Flemming3 reported that V addi-
6

7. tions are necessary to achieve strengths above 500 NITROGEN STRENGTHENING
350
MPa, and maximum strength steels were nitrogen-en-
Yield-Strength Increase, MPa
0.12% V
hanced to optimize the V(C,N) strengthening in the 300
ferrite. 0.09% V
250
TiN will tend to form during or soon after solidifi-
200
cation and may contribute to strengthening only if the 0.06% V
cooling rate is sufficient to keep the precipitates small11. 150
Generally, TiC can be effective for precipitation 100
strengthening, but the higher nitrogen levels of EAF 0.03% V
50
steels make it difficult to keep Ti in solution for subse-
quent TiC precipitation. 0
0 0.005 0.01 0.015 0.02 0.025
The presence of increased nitrogen in steel reduces
Nitrogen, wt. %
the relative strengthening effectiveness of both Ti and
Figure 9-- Effect of N content on yield-strength increase
Nb additions. The higher nitrogen levels of EAF steels
based on model of V-N microalloyed strip12.
typically used with thin-slab casting become a positive
factor when combined with V additions, increasing the tenite grains that occur from rolling deformation. These
effectiveness of the V addition. Strengthening coeffi- flattened grains then promote the formation of small
cients for N in V steels have been reported to be as ferrite grains in the final product.
high as 7 MPa (1 ksi) for each 10 ppm increment of N Nb-alloyed steels using these “controlled-rolling”
as long as the V:N ratio is greater than 4:112. Figures 8 (CR) processes were developed over 30 years ago. In
and 9 show the effects of nitrogen on strengthening the CR process, the temperature for the final rolling
vanadium steels as reported in the literature. passes must be held below the maximum temperature
at which recrystallization will not occur, the “recrys-
Grain-Size Considerations tallization stop temperature, Tnr”. In thin-slab process-
For maximum strength and ductility of polygonal fer- ing, the use of Nb to raise the Tnr temperature is com-
rite microstructures, it is advantageous to produce as plicated by the previously described problem of Nb
fine a ferrite grain size as possible. To achieve ferrite also inhibiting the necessary recrystallization of the as-
grain refinement in the final as-rolled product, the aus- cast microstructure.
tenite should be “conditioned” during the later stages Recrystallization controlled rolling (RCR) was de-
of the rolling process to provide the maximum amount veloped later. The RCR process relies on repeated
of grain-boundary surface area. One way this condi- interpass recrystallization to achieve a fine austenite
tioning can be accomplished is by finish rolling at tem- grain size, with controlled temperatures and reduction
peratures below the recrystallization stop temperature, schedules to achieve the desired results. The RCR roll-
thereby maintaining the flattened or “pancaked” aus- ing process involves finish-rolling temperatures more
typical for plain-carbon steels, while the CR process
NITROGEN STRENGTHENING requires lower finish rolling temperatures. These higher
300 rolling temperatures promote repeated recrystallization.
Precipitation Strengthening, MPa
Since niobium has been shown to raise the recrystalli-
0.22% C 0.15% C
250 zation stop temperature through precipitation and sol-
ute drag, V microalloying was naturally more com-
0.10% C patible with the RCR process. V does not inhibit the
200
austenite recrystallization required in the RCR process.
Both CR and RCR rolling generate an austenite
150
with a large amount of grain-boundary surface area
0.04% C
(austenite interfacial area), promoting transformation
100
to a small ferrite grain size. These small ferrite grains
contribute to both strength and to increased tough-
50 ness. The effect of the ferrite grain size on the final
0 0.005 0.01 0.015 0.02 0.025 0.03
properties of the steel is independent of the austenite
Nitrogen, wt. % conditioning process used to achieve that ferrite grain
Figure 8-- Precipitation strengthening of 0.12% V steels size. At an equivalent ferrite grain size, the contribu-
isothermally transformed at 650 deg. C / 500s15. tion of grain size to the final properties (strength and
7

8. toughness) will be comparable. Therefore, the test of microalloyed-strip steels produced on several differ-
the process is the final ferrite grain size and not neces- ent thin-slab cast mills12. The grain sizes for the higher
sarily the condition of the austenite that produced the yield-strength heats (>500 MPa) were in the 3 to 5µm
ferrite. For example, the presence of V and N can range, which is typical of the best found in any
change the ratio of given austenite grain size Dγ to the microalloyed-production steels.
transformed ferrite grain size Dα, as shown in Figure
10. A higher Dγ/ Dα ratio means that for a given aus- Strain-Aging Considerations
tenite grain size, the ferrite grain size will be smaller. Strain aging is often a concern in high-N steels.
In other words, the austenite grain size is not the only Strain aging is the increase in flow stress caused by
controlling issue in determining the final ferrite grain aging a strained material at slightly elevated tempera-
size13. tures. Nitrogen aging can occur from high ambient tem-
peratures upward, while carbon aging generally starts
TRANSFORMED FERRITE GRAIN SIZE at 150°C. With significant aging after cold deforma-
10
tion, the toughness and ductility of the steel can be
substantially degraded. The important consideration
Transformation Ratio, DγDα
5 here is that it takes “free” or uncombined interstitial
carbon or nitrogen to cause this aging phenomenon.
3 C-Mn Nitrogen that is tied up as a nitride in the steel will not
contribute to strain aging.
2 Strain aging can be evaluated as shown in Figure
0.1% V, 12. The strain-aging index, ∆Y, is the increase in flow
0.02% N stress after aging a pre-strained sample13. The pre-
strain was selected as 7.5%, with aging taking place at
100°C to insure that aging was only due to the pres-
1
1 10 100 1000
ence of free nitrogen.
Austenitic Grain Size, µm
STRAIN-AGING-INDEX TEST PROCEDURE
Figure 10 -- Effect of V and N on refining transformed ∆Y = CHANGE IN YIELD STRESS DUE TO STRAIN AGING
polygonal ferrite size15. ∆U = CHANGE IN ULTIMATE TENSILE STRENGTH DUE TO
STRAIN AGING
∆E = CHANGE IN ELONGATION DUE TO STRAIN AGING
The evaluation of production HSLA steels from
STRAIN-AGED LOWER
compact-strip processes has shown that an extremely YIELD EXTENSION
b
fine ferrite grain size can be produced with any of the ∆U
a
normal alloy designs. Figure 11 shows ferrite grain size ∆Y
A
as a function of gauge for a series of as-rolled V-N
Strain
B
FERRITE GRAIN SIZE VS. THICKNESS ∆E
< 500 M Pa > 500 M Pa INITIAL LOWER
YIELD
10 EXTENSION
(LÜDERS STRAIN)
9
Ferrite Grain Size, µm
8
7 PRE-STRAIN
6
5
4 Stress
3 Figure 12 -- A schematic drawing of aging-index (∆Y)
2 test procedure15.
1
0 Figure 13 shows the results of a strain aging evalu-
0.0 5.0 10.0 15.0 20.0 ation of sheet steels from thin-slab strip mills with EAF
Thickness, mm melting furnaces. There were from 5 to 10 different
Figure 11 -- Ferrite grain size of production samples of heats for each group. The results show that even with
V-N thin-slab strip steels of various yield strengths12. average nitrogen values of 150 ppm, there is no strain
8

9. STRAIN-AGING-INDEX RESULTS able without the need to roll below the recrystallization
Vanadium - 0.64% V, 0.009% N stop temperature.
Vanadium-Nitrogen - 0.067% V, 0.015% N
Carbon-Manganese - 0.00% V, 0.009% N • Processing requirements (reheating temperature,
40 rolling temperature, roll pass scheduling, and coiling
temperature) are comparable to processing C-Mn steels
Aging Index (∆ MPa
30 and do not require major modifications.
With the high degree of competitiveness in the cur-
(∆Y),
20 rent world market for strip steel, each steelmaking op-
eration will need to find the particular market niche
10 that suits the capabilities of that operation. High levels
of production in North America have confirmed that
0 thin-slab casting is particularly suited to producing hot-
V VN C-Mn rolled and microalloyed HSLA steels at competitive
-10 costs. The high compatibility of the V and the V-N
Steel Type alloying system to the production process requirements
Figure 13 -- Strain aging index results from V, V-N, and of thin-slab mills, as reviewed in this paper, has made
C-Mn steels16.. it the alloy of choice for most of these mills.
With over ten years of history and continuous im-
aging in the vanadium steels. The C-Mn steels, all of provement, vanadium-microalloyed HSLA strip steels
which were Al-killed, still had a significant strain-ag- from thin-slab casters are continuing to expand their
ing response. In as-rolled strip steels, V is much more usage into even more stringent quality applications.
effective in removing free nitrogen from solution than There is no known limit to prevent these mills from
is Al. Some of the V-N steels contained over 200 ppm supplying the majority of the market for HSLA hot-
nitrogen and still did not exhibit strain aging. rolled strip. The cost effectiveness of the V-N micro-
alloying system, along with the cost efficiencies of the
CONCLUSIONS thin-slab-casting operation, provide a highly competi-
Based on the previous discussion, vanadium has tive process that will be the standard for the future.
the capability of providing the stated microalloy char-
acteristics preferred for the thin-slab-cast, direct-charg- REFERENCES
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The Material Difference Is Value
For information on Stratcor® vanadium products, please contact:
STRATCOR, Inc.
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