precipitation strengthening but also for ferrite grain refinement. ..... into austenite where the growth is controlled by carbon diffusion in austenite while maintaining.
PRECIPITATION AND GRAIN REFINEMENT IN VANADIUM – CONTAINING STEELS Stanislaw Zajac Swedish Institute for Metals Research Drottning Kristinas väg 48, S-114 28 Stockholm, Sweden Tel: +46 8 440 4884 E-mail: stanislaw.zajac @simr.se
Abstract The present work has concentrated on the roles of vanadium, nitrogen and carbon in controlling the precipitation of V(C,N) in austenite and ferrite and their effects on; (i) grain refinement by promoting the formation of intragranular ferrite, and (ii) precipitation strengthening by interphase and random precipitation. The degree of precipitation strengthening of ferrite at a given vanadium content depends on the available quantities of carbon and nitrogen. It is concluded that nitrogen is a very reliable alloying element, increasing the yield strength of V-microalloyed steels by some 6 MPa for every 0.001% N, essentially independent of processing conditions. The role of carbon in precipitation strengthening is complicated. The present results have shown that the precipitation strengthening of V-microalloyed steels increases significantly with the total C-content, ~5.5MPa/0.01% C. The explanation is that the metastable equilibrium between ferrite and undercooled austenite greatly increases the solubility of carbon in ferrite, in the times available during transformation, thereby contributing to profuse nucleation of V(C,N) particles. This effect of carbon is particularly significant for medium carbon steels typically used for hot rolled bars and sections. The experimental results strongly indicate that vanadium can by effectively used not only for precipitation strengthening but also for ferrite grain refinement. It was shown that vanadium contributes to the formation of two types of intragranulary nucleated ferrite; polygonal (idiomorph) ferrite and acicular (sideplate) ferrite. Intragranular polygonal ferrite nucleates on VN particles that grow in austenite during isothermal holding or slow cooling throughout the austenite range. Acicular ferrite microstructure forms in V-microalloyed steels during isothermal transformation at lower temperatures (~450°C). The acicular ferrite microstructure was obtained in V-microalloyed steels containing high, medium or very low nitrogen levels. This suggests that vanadium on its own can promote the formation of the acicular ferrite microstructure. 1.
INTRODUCTION
Vanadium is best known as an eminent element for strong and easy controllable precipitation strengthening. The principal reason for this is the relatively large solubility product of its carbonitrides resulting in a lower solution temperature and a larger capacity to dissolve them at elevated temperatures. A special feature of V as compared to Nb is that its nitride is much less soluble than its carbide and this gives N a very important role in V-steels, especially in their precipitation
1
strengthening. In order to maximise the precipitation strengthening it is necessary to understand the roles of nitrogen and carbon in formation of high volume fractions of finely dispersed carbonitrides. Previous data strongly indicate that the role of nitrogen is clear1. It has been demonstrated that the yield strength increases almost linearly with increasing nitrogen content for given vanadium and carbon levels making N an eminent choice for strong and easily controllable precipitation strengthening2. Carbon content has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite. This was deduced from the fact of very restricted solubility of carbon in ferrite (which is normally supposed to be independent of the total carbon content in the steel). Most published literature does not suggest that differences in carbon contents in the range for structural steels (0.04-0.3%) should affect significantly the response of vanadium in these steels3. However, this viewpoint must be revised in the light of the later work2 which have shown that the precipitation strengthening increases significantly with total C-content of the steel. Increasing C-content delays the pearlite formation and thereby maintains for a longer time the higher content of solute C in ferrite corresponding to the austenite/ferrite equilibrium as compared to that of ferrite/cementite, allowing more nucleation of V(C,N) particles and accordingly a more dense precipitation. Recent studies at SIMR4 and literature data5-7 strongly suggest that vanadium can also by effectively used for ferrite grain refinement. There is also evidence that vanadium promotes the formation of acicular ferrite microstructure steels8. It was suggested that the VN particles which precipitated inside austenite grains during/or after hot rolling show strong potential for nucleation of intragranular ferrite. Although, vanadium does not readily precipitate in austenite, the precipitation process can be enhanced with increasing nitrogen in the steel or by plastic deformation (straininduced precipitation). In the case of precipitation in ferrite, the Baker-Nutting (B-N) orientation relationship is observed between the V(C,N) particles and matrix9. This fact may be very important for the nucleation of ferrite on cooling as the interfacial energy between ferrite and vanadium nitride is very low for the B-N relationship. In fact, intragranular ferrite idiomorphs were observed to nucleate at vanadium nitrides with the B-N orientation7. These observations suggest that intragranular ferrite can nucleate on VN and maintain coherent, low energy interfaces with respect to vanadium nitride. Thus, it was concluded that the main factor governing the formation of intragranular ferrite is the presence of VN precipitates in austenite which can develop coherent, low energy, interphase boundaries with ferrite10. The precise mechanism by which vanadium additions may enhance the nucleation rate of ferrite is not known7,10,11, but this new role of vanadium can be extremely important for thermo-mechanical processing as grain refinement leads to a significant increase in strength which is accompanied by a marked improvement in toughness. The aim of the present paper is to review the present knowledge on the roles of V, N and C in precipitation strengthening as well as on grain refinement by promoting the formation of intragranular ferrite. 2.
THERMODYNAMIC CONSIDERATIONS
2.1
Solubility and composition of V(C,N) in microalloyed steels
To achieve the desired metallurgical states, a detailed knowledge of the solubilities of the microalloy carbides and nitrides is required, together with a knowledge of their precipitation behaviour. An understanding of the role of V as well as C and N can be gained from the solubility data summarised in Figs 1-3 from a recent evaluation of the thermochemical parameters for multi-component systems of HSLA steels12.
2
Fig. 1 shows Thermo-Calc calculations of the equilibrium precipitation of V(C,N) in steels microalloyed with vanadium. Also shown in this figure is the mole fraction of nitrogen in carbonitrides at various temperatures and for various nitrogen contents from hyperstoichiometric levels to zero. These results show that vanadium starts to precipitate in austenite as almost pure nitride and the precipitation start temperature depends strongly on the nitrogen level. For 0.10%V and 0.03%N the precipitation starts at ~1110°C and the precipitation start temperature decreases to below 950°C at 0.003%N. When the nitrogen is about to be exhausted there is a gradual transition to form mixed carbo-nitrides, Fig. 1(b). Enhanced precipitation during the transformation is a result of the solubility drop of the vanadium carbo-nitrides associated with the transformation from austenite to ferrite at a given temperature. A significant feature of the solubility of VC in austenite is that this is considerably higher than the solubility of VN, suggesting that vanadium carbide will not precipitate in austenite.
Fig. 1 Precipitation of nitrides, nitrogen rich carbonitrides and carbides in 0.10% V steels at various nitrogen contents. The precipitation of carbo-nitrides in multiple microalloyed steels with V, Nb, Ti and Al is shown in Fig. 2(a), where the soluble components are expressed as a function of temperature. The mole fraction of Ti, Nb and V in M(C,N) is also included in this figure. As expected, the major part of the precipitate at the highest temperatures is TiN whilst further precipitation at lower temperatures is predominantly Nb(C,N). Thus, the primary precipitates to be formed in austenite are composite (Ti,Nb)-nitrides. Thermodynamic calculations imply that at the high temperature of 1200°C, primary (Ti,Nb,V)N particles contain ~20% of Nb and ~5% of V for multiple microalloyed steels. It is clear from this figure that the volume fraction of microalloy nitrides at high temperatures will be considerably larger in the Ti-Nb-V steel than in the single microalloyed steels. It may also be seen that AlN (with close packed hexagonal structure) has little or no solubility for other microalloying elements and starts to precipitate at approximately 1200°C in the presence of 0.035% Al, Fig. 2(a). Another important aspect of multiple microalloying is that the niobium or vanadium tied up as (Ti,Nb,V)N particles is not available for subsequent thermo-mechanical treatment. Fig. 2(b) shows that a significant fraction of added niobium may remain undissolved even at high reheating temperatures and cannot contribute to retardation of recrystallisation and/or precipitation strengthening. Vanadium, on the other hand, exhibits higher solubility and is therefore normally fully available for precipitation strengthening in ferrite.
3
Fig. 2 Precipitation of nitrides and nitrogen-rich carbonitrides in multiple microalloyed steel (a) and the mole fraction of Ti, Nb and V in M(C,N) (b). 2.2
Thermodynamic Driving Force for Precipitation
The precipitation process proceeds at a perceptible rate only if there is a driving force, that is, a free energy difference between the product and parent phases. This driving force enters the steady state nucleation rate in a central way and must be known with some accuracy if nucleation rates are to be calculated, or even estimated. The chemical driving force for nucleation of VN and VC in Vmicroalloyed steels, determined from the HSLA database are illustrated in Fig. 3. It can be seen that, as the temperature is decreased the driving force increases monotonically and changes slope after the austenite to ferrite transformation. The dominating effect of N on the driving force is clearly seen in Fig. 3(a).
(a)
(b)
Fig. 3 Chemical driving force, ∆Gm/RT, for precipitation of VC and VN in 0.12% V steel. For hard precipitates such as V(C,N) bypassing of dislocations is expected to occur by bowing between the particles (Orowan-mechanism) under all practical conditions. This means that the precipitation strengthening effect is determined by the interparticle spacing. In precipitation reactions, the decisive factor which minimises the interparticle spacing is the rate of nucleation since
4
it determines the particle density. Hence, the strong effect of N in increasing this driving force, Fig. 3(a), is the explanation of the well-known increase in strength by N in V-microalloyed steels. It should be noted, however, that the driving force can also be raised by V-addition but the effect per concentration unit is much less, Fig. 3(b). Moreover, V-additions increase alloying costs significantly whereas the usage of N at the present levels is free. 3.
PRECIPITATION OF V(C,N) IN AUSTENITE AND FERRITE
3.1
Kinetics of VN precipitation in austenite
The precipitation-temperature-time diagram for VN in undeformed (recrystallised) austenite is shown in Fig. 4 for the 0.12%V steel. The experimental results yield the well-known “C-curve” associated with the kinetics of VN precipitation. The curves indicate precipitated vanadium at the level of 0.001%, 0.005%, 0.01% and 0.02%. The location of the C-curve on time-temperature diagram is very important in regards to processing since precipitates will only form when the processing time-temperature profile passes through the C-curve and the effectiveness of VN in controlling transformation reactions is directly related to the amount and distribution of the precipitate which is subsequently related to the location of the C-curve. 1025 1000 975
Temperature, °C
0.001
0.1%C-0.12%V-0.0082%N 0.005
950
0.001
0.01
925 900 875 850
0.02 0.001
0.005
0.01
825 800 775 1.E+01
Precipitated V (wt%) 1.E+02
1.E+03
1.E+04
1.E+05
1.E+06
Time, s
Fig. 4 Precipiation-time-temperature diagram for VN in undeformed austenite. It is clearly seen from Fig. 4 that the precipitation process of VN in undeformed austenite is very sluggish. After holding for 1 hour at 850°C less than 10% of the equilibrium amount of V, which can precipitate at this temperature, was precipitated in VN. The maximum rate of precipitation of VN in austenite was evaluated to be in the range of 850-875°C. 3.2
Precipitation of V(C,N) in Ferrite
As regards strengthening, the effective V-carbonitrides are those formed in ferrite during the latest passage through the austenite-ferrite transformation. At equilibrium the circumstances are such that a certain, small portion of the vanadium in microalloyed steel should precipitate in austenite, especially if the contents of vanadium and nitrogen are high. However, the kinetics of V(C,N) formation in austenite are sluggish and for processing at finishing temperatures higher than 1000°C and for normal steel compositions virtually all vanadium will be available for precipitation in ferrite. 5
Precipitation of V-carbonitrides can occur randomly in ferrite in the wake of the migrating austeniteferrite (γ-α) boundary – general precipitation – or by interphase precipitation characterised by the development of sheets of particles parallel to the γ/α-interface formed repeatedly with regular spacing. Many investigations have shown that for compositions typical of structural steels the general precipitation takes place at lower temperatures, typically below 700oC, and the interphase precipitation at higher temperatures. It is now a well established fact, demonstrated in several electron microscopy studies1,2, that V(C,N)-particles are also formed in the pearlitic ferrite. Because of the lower transformation temperatures of pearlite this type of precipitation is usually finer. Both interphase and general precipitation have been observed. 3.2.1 Interphase precipitation Fig. 5 shows the typical morphology of interphase precipitation of V(C,N) in 0.04-0.10%C, 0.13%V steels. Already from its appearance one can conclude that such a microstructure is formed in sheets parallel to the γ/α-interface by repeated nucleation of particles as the transformation front moves through the austenite.
Fig. 5 Electron micrographs of the selected steels isothermaly transformed at 750°C for 500 s, (a) 0,0051%N (b) 0,0082%N (c) 0,0257%N (d) 0,0095%N, 0,04%C. The effect of transformation temperature (e) and N-content (f) on intersheet spacing of V(C,N) interphase precipitation in 0.040.10%C, 0.12%V, 0.005-0.025%N steels. At high transformation temperatures, ~800°C, the interphase precipitation consists of irregularly spaced, and often curved sheets of V(C,N)-particles1. With decreasing temperatures the incidence of curved rows of precipitates diminishes and the dominant mode is regularly spaced, planar sheets of particles. From about 700oC the interphase precipitation is commonly found to be incomplete, and random precipitation from supersaturated ferrite after the γ/α-transformation takes over progressively with decreasing temperature. A characteristic feature of interphase precipitation is that it becomes more refined at lower temperatures, Fig. 5(e). Fig. 5(f) shows that the intersheet spacing is affected considerably by the nitrogen content of the steel. As shown there, it is diminished to almost one third at 750oC on increasing the nitrogen content from 0.005 to 0.026%.
6
Raising the carbon content above 0.10% C, will similarly decrease the A3 temperature of the steel and suppress interphase precipitation. According to the present investigation, interphase precipitation was visible within proeutectoid ferrite or pearlitic ferrite only at transformation temperatures very close (within 40-50°C) to the eutectoid point. At lower transformation temperatures, formation of solute supersaturated proeutectoid ferrite and pearlitic ferrite occurs instead with random distributions of V(C,N). 3.2.2
The mechanism of interphase precipitation
The mechanism of interphase precipitation has been the subject of considerable discussions. The models that have been proposed to explain the phenomenon fall broadly into two categories; ledge mechanisms and models based on solute diffusion control. Honeycombe and coworkers were among the first to study interphase precipitation more profoundly13,14. They suggested that interphase particles form heterogeneously on γ/α-boundaries thereby pinning their migration normal to the boundary. Local breakaway leads to formation of mobile ledges. The ledges move sideways while the remaining part of the released boundary is stationary and enables repeated particle nucleation to occur, forming a new sheet. Hence, in this mechanism the intersheet spacing will be determined by the ledge height. From this follows one of the main drawbacks with the ledge mechanism, viz. to produce a credible explanation of the observed variation of the intersheet spacing with temperature, and steel composition, especially N, V and C. It is hard to see how these parameters should generate a corresponding variation of the ledge height. C
cV 0
cV N V
B
V
λ
V
V V
γ
V
α /γ
cV
α
cV
A
α /VCN
cV C
o B
A
x
Fig. 6 Left figure shows schematically how the γ/α-interface bows out, expands sideways and reaches eventually material with sufficient V for renewed precipitate nucleation to occur. The transfer of V boundary diffusion to the lower precipitate row is indicated. The right figure shows the V-content profile in a cross section, see eq(1). Among the models based on diffusion control the solute-depletion model due to Roberts15 is the most prominent and promising one and has been recently further developed by Lagneborg and Zajac16. The quantitative description of the solute-depletion model considers a ferrite grain growing into austenite where the growth is controlled by carbon diffusion in austenite while maintaining local equilibrium at the interface, Fig. 6. At a given point of this growth the interplay is analysed between the nucleation of V(C,N)-particles in the γ/α-interface, the accompanying growth of Vdepleted zones around the precipitates, and the continued migration of the γ/α-boundary away from the precipitate sheet. The following relationship for the intersheet spacing, λ, has been suggested;
7
2 D b S c α / γ * − cVα / VCN λ = V V 0 α /γ* K 1 cV − cV
(1)
where S is the distance from the point of α-nucleation to the location of the γ/α-interface at time t, α / VCN DVb the boundary diffusion of V, cV the V-content of α in the α/VCN interface, cVα / γ * the critical V-content in the γ/α interface for nucleation of interphase precipitation, K1 is the proportionality factor for the ferrite growth. It was shown that volume diffusion of V cannot explain the observed intersheet spacings and that a faster diffusion process is required. Hence, a new mechanism was put forward where the transport of V occurs by boundary diffusion in the moving α/γ-interfaces. The model exhibits good agreement with observed values of intersheet spacing and their dependences on C-, V-, N-contents and temperature, as illustrated in Fig. 7. The computations shown in Fig. 7 also predict that the intersheet spacing around 700oC will fall below the approximate size of observed precipitates. Physically, this implies that the growth rate of the γ/α-interphase exceeds that of the V-depleted zone. Hence, the γ/α-boundary escapes, leaving the ferrite in its wake supersaturated with respect to V(C,N).
300
Intersheet spacing, nm
250
0,10C-0,06V-0,010N, calculated 0,10C-0,06V-0,010N, experimental 0,10C-0,12V-0,010N, calculated 0,10C-0,12V-0,010N, experimental 0,04C-0,12V-0,010N, calculated 0,04C-0,12V-0,010N, experimental
200
150
100
50
0 680
700
720
740
760
780
800
Temperature, °C
Fig. 7 Experimentally measured and computed intersheet spacings in interphase precipitation of a 0.06-0,13%V steel as a function of the transformation temperature and carbon and nitrogen contents. The model predicts also that the spacing is directly proportional to the ferrite growth, or to the degree of transformation. In fact, the model predicts that in the early stages of transformation the ferrite growth will be too rapid for V(C,N)-nucleation to take place and it is only later, when the growth rate has declined sufficiently, that the condition of interphase nucleation will be fulfilled. In the first part of the γ/α-transformation the ferrite will therefore be supersaturated with respect to V(C,N) directly behind the moving γ/α-interface and hence general precipitation will occur. Smith and Dunne17 made exactly this statement based on their microscopical observations and without the assistance of a model. Beside this, however, no attention seems to have been paid in the literature to this essential feature of interphase precipitation. The general statement often made, that a large variation of precipitation modes are observed, is probably largely a reflection of this specific characteristic of the phenomenon.
8
3.2.3
General precipitation
For compositions typical of V-microalloyed steels, 0.10%C-0.10%V, general precipitation of V(C,N) occurs in a temperature range from about 700oC and below. As already discussed in the previous section this transition from interphase to general precipitation can be predicted by the solute-depletion model. Randomly distributed V(C,N) particles are most common in V-steels transformed at 550-650°C at all carbon contents. The fact that these particles had formed in the ferrite after transformation and not on the inter-phase boundary during transformation is evident both from their random dispersion and from the existence locally of several different variants of the B-N orientation relationship9. Experimentally it is well established that VN has a considerably lower solubility and much larger chemical driving force for formation than VC, both in ferrite and austenite. This larger chemical driving force makes N-rich V(C,N) the preferred precipitation as long as there is sufficient nitrogen in the matrix, as has been demonstrated both by thermodynamic analysis, Figs 1 and 3 and by experiment18. It is only when the nitrogen content falls below about 0.005% that the initial V(C,N) starts to increase its C-content, Fig. 8.
Fig. 8 Thermodynamic calculations of the variation in the composition of VCxNy with content of N remaining in solution during precipitation in ferrite18.
Fig. 9 Growth of V(C,N)-precipitates after transformation at 650°C (a) as a function of the N-content of the steel, and (b) as a function of holding time.
A technically very important finding is that the precipitate size diminishes considerably with increasing nitrogen content in the steel, as is shown in Fig. 9. This is accompanied by a concurrent increase of precipitate density, although that is more difficult to measure quantitatively. These effects must be accounted for by an increased nucleation frequency as a result of the larger chemical driving force for precipitation of nitrogen-rich V(C,N). For the temperature and holding times in the experiments of Fig. 9 supersaturation with respect to V(C,N) still remains. Therefore the observed difference in precipitate growth between high and low-N steels cannot be explained by differences in coalescence. Instead, it must be interpreted as resulting from the denser particle nucleation in the high-N material, thereby producing an earlier soft impingement of V-denuded zones and so slowing down the precipitate growth. In numerous investigations of V-microalloyed steels it has been shown 9
that the observed precipitation strengthening originates seemingly only from N-rich V(C,N), despite the fact that there is abundant V to combine with the carbon dissolved in ferrite. The normal explanation of this behaviour has been that when the first formed N-rich precipitates have consumed practically all nitrogen the chemical driving force for formation of C-rich V(C,N) is too small for profuse precipitation to occur and hence no added strength is observed. However, this issue is complex, and recently it has been shown that under certain conditions the C-content in the steel can add significantly to precipitation strengthening, as shown below. The recent results of detailed measurements of particle sizes on steels having different carbon and nitrogen contents are summarised in Fig. 10, after isothermal transformation at 650°C. As indicated in this figure both nitrogen and carbon cause a decrease in the particle size and at the same time an increase in volume fraction, so giving the lowest planar spacing at a given transformation temperature. For steel having 0.005% nitrogen and 0.1% carbon the particles were relatively sparse with an average diameter of ~ 11 nm. They became smaller and more dense with increasing nitrogen content, reaching ~ 6 nm at 0.022% nitrogen. An increase in carbon content from 0.10% to 0.22% causes significant refinement (about a factor of two) of the vanadium-rich particles at all nitrogen levels. A higher density of smaller V(C,N) suggests that carbon play an important role in precipitation of V(C,N). This agrees with the suggestion that metastable carbon content increases the supersaturation levels achieved in ferrite and the intensity of V(C,N) precipitation is higher. 12
PARTICLE DIAMETER, nm
(d)
Isoth. trans. 650°C/500s
10
8
0.10C-0.12V
6
4
0.22C-0.12V 2
(high density)
(low density)
0 0
0.005
0.01
0.015
0.02
0.025
0.03
wt% N
Fig. 10 Electron micrographs showing V(C,N) precipitation after isothermal transformation at 650°C in 0,12%V-0,013%N steels with different C-contents; (a) 0.04% C, (b) 0.10% C and (c) 0.22% C , and the effect of N and C on the size and density of V(C,N)-precipitates during isothermal transformation at 650°C for 500s (d). 4.
PRECIPITATION STRENGRHENING
As was stated above vanadium is one of the most widely used precipitation strengthening additives in microalloyed steels. A modest addition of 0.10%V can bring about a strength increase beyond 200 MPa, and in special cases even up to 300 MPa, Figs. 11 and 1219. These figures show another essential feature of V-steels, viz. that nitrogen in small contents adds significantly to precipitation strengthening. Similarly, enhanced cooling through the γ/α-transformation and afterwards augments the particle hardening, Fig. 12. The results of Fig. 11 show also that normalising at 950oC does not allow all V to go into solution completely and hence produces only modest strengthening. 10
Fig. 11 Effect of processing method on the precipitation strengthening derived from V(C,N) for 0,12%C-0,35%Si-1,35%Mn0,09%V-0,02%Al steel19.
Fig. 12 Influence of cooling rate on the strength contribution from V(C,N) precipitates. Base composition as for Fig. 1118.
Fig. 13 Effect of N, V and transformation temperature on the precipitation strengthening in 0,1%C-V-N steels after isothermal ageing at different temperatures for 500s.
The technically very important effect of N on the strengthening of V-steels is illustrated in Fig. 13 for different ageing temperatures. This strong effect of N has sometimes been interpreted such that only VN forms as precipitates, and that the remaining V does not combine with C, possibly due to the larger solubility and lower chemical driving force for VC. A physically more correct account would be as follows. In the present material with hard non-penetrable V(C,N)-precipitates the particle strengthening occurs by the Orowan mechanism – bowing of dislocations between particles – and in that case the decisive parameter is the interparticle spacing in the slip plane. In turn, that is determined by the density of the precipitates. This density is controlled by the nucleation frequency and the number of nuclei it creates until the supersaturation has diminished so that nucleation dies. The essential parameter governing the variation in nucleation is the chemical driving force for V(C,N)-precipitation which, as was shown above, depends on the available quantities of both C and
11
N. It is true, however, that it varies relatively strongly with the N-content, but carbon may also significantly influence the chemical driving force.
Fig. 14 Deduced values of precipitation strengthening for isothermally transformed V-steels (650°C/500s). (a) as a function of N-content (b) as a function of C-content. Recent investigations2,20 have demonstrated very clearly that the precipitation strengthening of Vsteels increases significantly with the C-content of the steels. Plots of the observed strengthening ∆Rp against the C-content, and for comparison also N-content, are given in Fig. 14. These results show that carbon raises ∆Rp by ~ 5.5MPa for every 0.01%C in the steel, whereas nitrogen raises it by ~ 6 MPa for every 0.001%N. This important effect of carbon has not previously been noticed. The reason is presumably that there has always been an uncertainty in the procedure of evaluating the precipitation strengthening as to how the contribution of pearlite should be deduced and deducted from the measured yield strength. This procedure was improved in the present case, and furthermore nano-hardness measurements were performed in order to unequivocally demonstrate the precipitation hardening in the ferrite, Fig. 15. These results demonstrate very clearly that the precipitation strengthening of V-steels increases significantly with the C-content of the steels. The role of carbon in precipitation strengthening of V-steels is discussed below.
NANO-HARDNESS OF FERRITE, GPa
5
0.12% V
4.5
0.022%N 4
0.005%N
0.025%N
3.5 3
0.005%N 2.5
Isothermal transformation 650°C/500s
2 0
0.05
0.1
0.15
0.2
0.25
0.3
Fig. 15 The effect of carbon and nitrogen contents on the nano-hardness of ferrite in 0.12%V.
wt% C
12
4.1
The Role of Carbon in Precipitation Strengthening
Carbon content has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite. This was deduced from the fact of very restricted solubility of carbon in ferrite (which is normally supposed to be independent of the total carbon content in the steel). Most published literature does not suggest that differences in carbon contents in the range for structural steels (0.04-0.3%) should affect significantly the response of vanadium in these steels3. However, it was suggested recently that the effective carbon for precipitation in ferrite may be much greater in the times available during phase transformation20. Fig 16(a) shows the considerable difference in solubility of C in ferrite when governed by the γ/α and cementite/α equilibria below A1. The solubility at 600oC is 5 times larger in γ/α than cementite/α. This metastable equilibrium between ferrite and undercooled austenite can greatly increase the solubility of carbon in ferrite thereby contributing to profuse nucleation of V(C,N) particles. The driving forces for nucleation of V(C,N) calculated for two different levels of dissolved carbon content equivalent to the equilibration ferrite+pearlite or metastable ferrite+austenite are shown in Fig. 16(b). In the extreme case of 250 ppm carbon in ferrite the driving force for nucleation increases by about 25-30% in comparison with the equilibrium content at 650°C. This demonstrates that the nucleation of V(C,N) is indeed a function of the carbon content dissolved in ferrite. In reality the carbon content in ferrite should lie somewhere between these values depending on transformation kinetics, mainly phase boundary mobility. The total C-content affects the kinetics of the transformation of γ to α and an increase in total carbon content displaces the pearlite formation to longer times as was shown in Fig. 16(c). The implication of this is that a larger chemical driving force for V(C,N)-precipitation will remain longer and hence promote profuse nucleation. Consequently, denser V(C,N) precipitation nucleates in ferrite, as was shown in Fig. 10 for 0.22%C steels. This role of carbon in precipitation strengthening is particularly significant in steels with higher carbon contents where the metastable condition is more extreme and prolonged.
DRIVING_FORCE for V(C,N)
6
250 ppm carbon in ferrite
5.5
5
4.5
equilibrum carbon content in ferrite 4
3.5 0
50
100
150
200
250
300
Nitrogen content in ferrite, ppm
(a)
(b)
(c)
Fig. 16 (a) Solvus lines for C in ferrite in equilibrium with cementite and austenite. (b) Chemical driving force, ∆Gm/RT, for precipitation of VC and VN in 0.12% V steel. (c) Effect of equilibrium and metastable carbon content in ferrite on driving force for precipitation of V(C,N) at 650°C.
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In practice the metastable α-γ equilibrium solubilities may not be so large as shown in Fig. 16(a) since diffusion of carbon in austenite occurs away from the interface tending to lower the concentration there. This will be accentuated if the interface migration is hindered by some effect such as solute drag of substitional elements. In the extreme case where boundary mobility is very low and full equilibration of carbon can take place both in ferrite and austenite the new metastable carbon levels can be calculated with the aid of Thermo-Calc (activities of carbon in ferrite and austenite are equal). Fig. 17 shows such calculation for two different carbon contents and demonstrates that the level of carbon dissolved in ferrite which governs the nucleation of V(C,N) is indeed a function of the total carbon content. In reality the carbon content in ferrite should lie somewhere between these values and the limit given in Fig. 16(a) (i.e. ~ 0.024 %C at 650°C) depending on phase boundary mobility. The above situation continues until pearlite is nucleated at which point a new equilibrium with a reduced carbon activity is established. So fast is the diffusion of carbon in ferrite that the transition can take place in less than one second for typical microstructures. The great reduction in carbon activity at this stage is likely to virtually eliminate further nucleation of V(C,N) although some continued growth of existing particles may be expected. The results in Fig. 17(a) also predict that the metastable carbon content in ferrite depends on the degree of transformation. In the early stage of transformation the carbon content is only slightly higher than the equilibrium level. As degree of transformation increases, ferrite is progressively enriched with carbon, from the shrinking austenite, so producing more nucleation of V(C,N) particles and accordingly a more dense precipitation. This was confirmed from micro-hardness measurements of partially transformed samples of low N steels, as shown in Fig. 17(b). In these ways the positive influence of total carbon content on the precipitation strengthening contribution can be understood. 250
0.025
MICRO-HARDNESS OF FERRITE, HV15g
0.22%C 0.10%C
wt% CARBON
0.02
0.015
0.01
ferrite 0.005
240
0.22C-0.12V-0.022N
230
0.22C-0.12V-0.0015N
IT at 650°C
220 210 200 190 180 170 160 150
0 0
10
20
30
40
50
60
70
% Transformed ferrite
(a)
80
90
100
0
50
100 150 200 250 300 350 400 450 500 550
TIME at 650°C, sec
(b)
Fig. 17 Values of the metastable carbon content in ferrite as a function of volume fraction of transformed ferrite at 650°C (a) and micro-hardness of ferrite as a function of transformation time (transformed ferrite) in high N and low N steels (b).
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5.
GRAIN REFINEMENT
5.1
Formation of Intragranular Ferrite in V-Steels
The effect of V on the intragranular ferrite formation was examined on specimens isothermally treated at 850°C for 1 hour and partially transformed at the temperatures 700-450°C followed by gas quenching to room temperature. The resulting microstructures of the 0.22%C-0.12%V-0.013%N steel are shown in Fig. 18. High transformation temperatures resulted in largely reconstructive transformation at the prior austenite grain boundaries. The microstructure produced on a partial transformation at 700°C and gas quenching is shown in Fig. 18(a). It consisted of coarse grain boundary ferrite. Within the prior austenite grains a mixture of acicular ferrite with a second phase was formed as a result of the secondary transformation at a lower temperature during gas quenching. The microstructure obtained in his steel after a partial transformation at 650°C followed by gas quenching is shown in Fig. 18(b). Comparing to the specimen transformed at 700°C there was less grain boundary ferrite at the prior austenite grain boundaries, but a high density of intragranulary nucleated polygonal ferrite inside the prior austenite grains. Some of the intergranulary nucleated ferrite crystals were elongated with sharp edges which is probably characteristic of the faster cooling after the interrupted transformation. Quenching from 600°C produced a thinner band of grain boundary ferrite and higher number of intragranulary nucleated ferrite crystals, Fig. 18(c). There is also a change in the morphology of the intragranular ferrite which become less polygonal and more sideplate. However, it is not clear whether this plate morphology of ferrite was associated with reconstructive growth mechanism during isothermal transformation at this temperature or whether it formed at lower temperatures after the gas quench. There were also some ragged ferrite plates growing from the grain boundary ferrite allotriomorphs. a
b
d
e
c
(a) (b) (c) (d) (e)
700°C 650°C 600°C 550°C 450°C
Fig. 18 Microstructures of 0.22%C-0.12%V-0.013%N steel austenitised at 1150°C, cooled at a rate of 2°C/s to 850°C, held at this temperature for 1 hour and partially isothermally transformed between 700-450°C prior to gas quenching to room temperature.
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Fig. 18(d) shows the microstructure resulting from a partial isothermal transformation at 550°C followed by gas quenching. There was very little grain boundary nucleated allotriomorphs which a very thin skeleton of ferrite at the prior austenite boundaries. Within the prior austenite grains a high density of small ferrite crystals was nucleated. A large proportion of this intragranular ferrite was in the form of thin elongated plates, some of which were observed to have grown from particles. The microstructure produced on isothermal transformation at 450°C is shown in Fig. 18(e). Almost no grain boundary nucleated allotriomorphic ferrite is seen. The general matrix microstructure formed at this temperature was very fine acicular ferrite. 5.2
Number Density of Grain Boundary and Intragranular Ferrite
Metallographic examination of these specimens was carried out to determine the amounts of grain boundary ferrite, intragranular polygonal ferrite and intragranular acicular ferrite as a function of the transformation temperature. Measurements were made of the number density of ferrite grains to quantitatively describe the refining effect of intragranular ferrite formation. Fig. 19 summarises the effect of isothermal transformation temperature on the number density of grain boundary ferrite and intragranular ferrite. It can be seen that the grain boundary ferrite started to form directly below the transformation start temperature and the number density of grain boundary ferrite grains increased slightly with decreasing the transformation temperature. More complex behaviour was observed for intragranular ferrite. There was no intragranulary nucleated ferrite at higher transformation temperatures, above about 650°C. The temperature range where the large density of intragranular polygonal ferrite was formed was narrow, between 650-600°C. On further lowering the isothermal transformation temperature, below 600°C, there was a tendency for the intragranular ferrite to change morphology and grow as ferrite sideplates. The number density of intragranular ferrite grains at 550°C was higher than that for 600°C. The number density and morphology of the ferrite grains changed considerably after isothermal transformation at 450°C. Cooling to 450°C led to the formation of exclusively acicular ferrite structure with a more than 4 times higher density of acicular grains. The tendency for grain boundary ferrite formation was reduced and it was difficult to estimate the number density of grain boundary ferrite grains as they have a similar acicular morphology.
acicular ferrite
Intragranular ferrite
2
100 or intragranular ferrite/100um
Number of grain boundary ferrite/100um
120
Grain boundary ferrite 80
60
40
20
0 400
500
600 Temperature, °C
700
800
Fig. 19 Number density of grain boundary ferrite and intragranular polygonal and acicular ferrite grains as a function of transformation temperature.
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5.3
Effect of Vanadium on the Formation of Acicular Ferrite.
The experimental results described above indicate that intragranular acicular ferrite was the dominant microstructure after isothermal transformation at 450°C. This treatment was repeated for steels with low and high nitrogen contents in order to verify the role of N as well as VN in the formation of acicular ferrite in V-steels. The resulting microstructures are shown in Fig. 20. It is clearly seen that a similar acicular ferrite structure was obtained in the V-microalloyed steel with a very low N content of 0.0015%N and high N content of 0.025%N. In the low N steel no VN precipitates were expected to grow in austenite. This result confirms the suggestions of He and Edmonds8 that there exists an effect of vanadium on the formation of acicular ferrite microstructure even in low nitrogen steels. 0.22C-0.12V-0.0015N
0.1C-0.12V-0.025N
Fig. 20 Microstructure of low N-steel, and high N-steel isothermally transformed at 450°C. 5.4
Analysis of Nucleation Sites
Nucletion of intragranular ferrite and the morphology of nucleating precipitates/inclusions was examined in several of V-containing steels using both scanning and transmission electron microscopy. Fig. 21 shows a micrograph of the 0.10%C-0.12%V-0.025%N steel, which was water quenched from 630°C after ~5% transformation. In the intragranular regions, a number of inclusions exhibited ferrite growing with different morphologies. There were loops, blocks and caps of ferrite coating part of the inclusion surfaces as well as wedges and ferrite growing with more of a plate type of morphology. The formation of very long plates from individual inclusions was also observed. A number of spherical non-metallic inclusions were present in the investigated steels. The majority of these inclusions were identified by EDS as manganese sulphides, Fig. 21. Other inclusions which were identified were oxides, mainly silicon-manganese oxides. Vanadium-rich particles, presumably nitrides were always visible in connection with inclusions. The fact that VN particles had formed in the austenite prior to transformation and nucleated ferrite crystals is clearly evident from SEM micrographs.
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MnS
VN
α
Fig. 21 Intragranular ferrite nucleus in 0.10%C-0.12%V-0.025%N steel, water quenched from 630°C after ~5% transformation. 6.
SUMMARY
The present work has concentrated on two strengthening mechanisms in V-microalloyed steels: (i) grain refinement by promoting the formation of intragranular ferrite, and (ii) the role of nitrogen and carbon in precipitation strengthening by interphase and random precipitation of V(C,N) in ferrite. The experimental results strongly indicate that vanadium can by effectively used not only for precipitation strengthening but also for ferrite grain refinement. It was shown that vanadium contributes to the formation of two types of intragranulary nucleated ferrite; polygonal (idiomorph) ferrite and acicular (sideplate) ferrite. Intragranular polygonal ferrite nucleates on VN particles which grow in austenite during isothermal holding or slow cooling throughout the austenite range. The intragranular polygonal ferrite forms in the narrow temperature range, between 650-600°C for the investigated compositions. Acicular ferrite microstructure forms in V-microalloyed steels during isothermal transformation at lower temperatures (~450°). The acicular ferrite microstructure was obtained in V-microalloyed steels containing high, medium or very low nitrogen levels. This suggests that vanadium on its own can promote the formation of the acicular ferrite microstructure. Vanadium is an effective and easy controllable precipitation strengthening element. The degree of precipitation strengthening of ferrite for a given vanadium content depends on the available quantities of carbon and nitrogen. It was confirmed that nitrogen is a very reliable alloying element, increasing the yield strength of V-microalloyed steels by some 6 MPa for every 0.001% N, essentially independent of processing conditions. Experimentally it is demonstrated that the V(C,N)precipitation becomes denser and finer with increasing N-content. Carbon content, on the other hand, has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite because of the very small carbon content in solution in ferrite at equilibrium. The present results have shown that the precipitation strengthening of V-microalloyed steels increases significantly with the total C-content, ~5.5MPa/0.01% C. The explanation is that the C-content of the steel delays the pearlite (ferrite+cementite) formation and thereby maintains the higher, metastable C-content in ferrite given by the austenite/ferrite equilibrium for a longer time. This effect of carbon is particularly significant for medium carbon steels typically used for hot rolled bars and sections.
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7.
ACKNOWLEDGEMENTS
The present endeavour has been supported financially by the U.S. Vanadium Corporation /Stratcor/ and the Vanadium International Technical Comittee /VANITEC/. It is a pleasure to acknowledge the stimulating and challenging discussions with Michael Korchynsky of U.S. Vanadium and Peter Mitchell of VANITEC. Thanks are also due to Bevis Hutchinson and Rune Lagneborg for valuable discussions. REFERENCES 1.
S. Zajac, T. Siwecki and M. Korchynsky: “Importance of Nitrogen for Precipitation Phenomena in V-Microalloyed Steels”, in Proc. Conf. Low Carbon Steels for the 90`s, (eds. R Asfahani and G. Tither), ASM/TSM, Pittsburgh, USA, 1993, pp.139-150.
2.
R. Lagneborg, T.Siwecki, S. Zajac and B. Hutchinson, “The Role of Vanadium in Microalloyed Steels”, Scandinavian Journal of Metallurgy, Vol 28, No. 5, October 1999, pp.1241.
3.
W. Roberts, “Recent Innovations in Alloy Design and Processing of Microalloyed Steels”, 1983 Int. Conf. on Technology and Applications of High Strength Low Alloy (HSLA) Steels, Philadelphia, PA, Oct. 3-6, 1983, pp. 67-84.
4.
S. Zajac, T. Siwecki, B. Hutchinson, L-E. Svensson and M. Attlegård, ”The Influence of Plate Production Processing Route, Heat Input and Nitrogen on the HAZ Toughness in Ti-V Microalloyed Steel”, Swedish Institute for Metals Research, Report IM-2764, (1991).
5.
T. Kimura, A. Ohmori, F. Kawabata, and K. Amano ”Ferrite Grain Refinement through Intragranular Ferrite Transformation VN Precipitates in TMCP of HSLA Steel”. Thermec 97, (1997) pp.645-651.
6.
T. Kimura, F. Kawabata, K. Amano, A. Ohmori, M. Okatsu, K. Uchida and T. Ishii, “Heavy Gauge H-Shapes with Excellent Seismic-Resistance for Building Structures Produced by the Third Generation TMCP”, CAMP-ISIJ, (1999) pp.165-171.
7.
F. Ishikawa and T. Takahashi, ”The Formation of Intragranular Ferrite Plates in Medium-Carbon Steels for Hot-Forging and its Effect on the Toughness”, ISIJ International, (1995) pp.1128-1133.
8.
K. He and D.V. Edmonds, “The Formation of Acicular Ferrite in C-Mn-V Steels”, 19th ASM Heat Treating Society Conf., Cincinnati, OH, (1999) pp.519-525.
9.
T. N. Baker, "Precipitation and Transformation in Steels", Acta Metall., 23, (1973, p. 261.
10.
M. Enomoto, “The Mechanisms of Ferrite Nucleation at Intragranular Inclusion in Steels”, International Conference on Thermomechanical Processing of Steel & Other Material, Edited by T. Chandra and T. Sakai, The Minerals, Metals & Materials Society, (1997) pp.427-433.
11.
F. Ishikawa, T. Takahashi and T. Ochi, “Mechanism of Intragranular Ferrite Nucleation”, Solid-Solid Phase Transformations, Eds. W.C. Johnson, J.M. Howe, D.E. Laughlin and W.A. Soffa, The Minerals, Metals & Materials Society, (1994) pp.171-176.
12.
S. Zajac, “Thermodynamic Model for the Precipitation of Carbonitrides in Microalloyed Steels”, Swedish Institute for Metals Research, Report IM-3566, Jan. 1998.
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13.
A. T. Davenport and R. W. K. Honeycombe, “Precipitation of Carbides at Austenite/Ferrite Boundaries in Alloy Steels”, Proc. Roy, Soc. London 322 (1971) pp. 191-205.
14.
R. W. K. Honeycombe, “Transformation from Austenite in Alloy Steels”, Metall.Trans.A, 7A (1976) pp. 915-936.
15.
W. Roberts, Hot Deformation Studies on a V-Microalloyed Steel, Swedish Institute for Metals Research, Report IM-1333, 1978.
16.
R. Lagneborg and S. Zajac, “A Model for Interphase Precipitation in V-Microalloyed Structural Steels”, Metall. and Mat. Trans. Vol. 32A, January 2001, pp.39-50.
17.
R. M. Smith and D. P. Dunne, Structural Aspects of Alloy Carbonitride Precipitation in Microalloyed Steels, Materials Forum 11 (1988) pp.166-181.
18.
W. Roberts, A. Sandberg, The Composition of V(C, N) as Precipitated in HSLA Steels Microalloyed with Vanadium, Swedish Institute for Metals Research, Report IM-1489 Oct. 1980.
19.
W. Roberts, A. Sandberg, and T. Siwecki, Precipitation of V(C,N) in HSLA Steels Microalloyed with V, Proc. Conf. Vanadium Steels, Krakow, Oct. 1980, Vanitec, pp.D1-D12.
20.
S. Zajac, T.Siwecki, W. B. Hutchinson, R. Lagneborg., “Strengthening Mechanisms in Vanadium Microalloyed Steels Intended for Long Products”, ISIJ International, 38 (1998) pp.1130-1139.
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