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The influence of manganese and silicon on the precipitation of vanadium carbide in steel
Materials Science and Engineering, AI 11 (1989) 189-199
189
The Influence of Manganese and Silicon on the Precipitation of Vanadium Carbide in Steel H. S. UBHI* and T. N. B A K E R
Division of Metallurgy and Engineering Materials, Universityof Strathclyde, Glasgow, G I 1XN (U.K.) (Received November 2, 1987; in revised form October 31, 1988)
Abstract
Four vacuum-melted steels with a nominal composition of Fe-O.12wt. %C-0.46wt. %V contained additions of 1.5 wt. % Mn (steel B), 0.3 wt. % Si (steel C) and 1.5 wt. % Mn and 0.3 wt. % Si (steel D). Steel A, with no manganese or silicon additions, was the control steel. After hot rolling, some of the material was drawn to wire, prior to solution treating at 1250 °C followed by aging for different times at 650 °C. Resistivity, hardness and metallography were used to follow the precipitation kinetics of vanadium carbide. Alloy carbide precipitation at 650 °C was found to start simultaneously with austenite transformation and occurred only in the ferrite phase. Silicon increased the rate of precipitation by approximately four times, but manganese, other than affecting the rate of austenite transformation, had little influence on the alloy precipitation kinetics. The combined addition of manganese and silicon resulted in an intermediate rate of transformation which influenced the precipitation kinetics of vanadium carbide.
tion strengthening, further tempering may be required [6, 7]. In order, therefore, to optimize the benefit from precipitation strengthening, a knowledge of the precipitation kinetics is of vital importance. However, few data on this topic are available; hence the object of this work is to investigate the kinetics of vanadium carbide precipitation in 0.12wt.%C-0.46wt.%V steels as well as the effects of manganese and silicon additions when added singly and together. This has been done by exploiting the extreme sensitivity of electrical resistivity changes in metals due to processes such as precipitation and dislocation annealing. In addition, transmission electron microscopy has been used to give structural information as a complement to the electrical resistivity data. This study is confined to precipitation during isothermal aging and is an extension of previous studies on the base steel [8]. Vacuum-melted microalloyed steels have been used to study the influence of manganese and silicon on the rate of isothermal precipitation of vanadium carbide in the ferrite phase. The four steels (Table 1) used had a nominal composition
I. Introduction
TABLE 1
High strengths in low alloy steels are achieved primarily by a combination of grain refinement and precipitation hardening [1]. Whereas the former can be produced by using stringent rolling procedures [2], the latter usually occurs both during and after rolling [3-5]. However, it is evident that, to reach the full potential of precipita-
Steel
*Present address: Department of Mechanical Engineering, University of Nairobi, Kenya. 0921-5093/89/$3.50
A B C D
Alloy compositions
Amount (wt.%) offollowing elements V
C
Si
Mn
0.46 0.46 0.46 0.45
0.14 0.14 0.12 0.12
0.01 0.01 0.28 0.30
0.01 1.5 0.01 1.5
The steels also contained 0.004 wt.% P, 0.003 wt.% S, less than 0.01 wt.% Cr, less than 0.005 wt.% Mo, less than 0.01 wt.% Ni, less than 0.005 wt.% AI, less than 0.001 wt.% N and less than 0.006 wt.% O, with the balance iron.
© Elsevier Sequoia/Printed in The Netherlands
190
of Fe-0.12wt.%C-0.46wt.%V. Manganese at the 1.5 wt.% and silicon at the 0.3 wt.% levels were present singly or in combination. The carbon-totransition-metal ratio was chosen to be approximately equivalent to the stoichiometric VC ratio.
1.0 ~
'
~
oA OB
oC +D 0.8
>,
2. Experimental details The vacuum-melted steels, the compositions of which are given in Table 1, were supplied in the form of hot-rolled 32 mm 2 square bars. The bars were cut into 150 mm lengths, hot rolled to rod of 10 mm diameter and then cold drawn to wire of 3 mm diameter. Specimens of 100 mm length were cut from this wire and preheat treated in encapsulated silica tubes for 5 h at 1200 °C followed by air cooling. They were then re-encapsulated, solution treated for 15 min at 1250 °C followed by isothermal aging at 650 °C in a tin bath for specified aging times. After this time lapse, the specimens were quickly lifted out of the tin bath, the silica tube was broken open and the specimen was allowed to drop into a quenching tank containing a medium of 4% aqueous brine. This operation was completed within 15 s. The voltage drop across the aged specimens was measured in liquid nitrogen, using the classical four-point method. A current of 0.5 A from a Hewlett-Packard G177B source was passed through the specimen in series with a 0.20 f2 standard resistance held in a constant-temperature bath. The potential drop across both the specimen and the standard resistance was measured with a Solartron A210 digital voltmeter. Resistivity values were obtained to an accuracy of + 2%. Both thin foil and carbon extraction replicas were prepared by standard techniques and examined using a Philips EM400T transmission electron microscope. 3. Results Normalized resistivity p , vs. isothermal aging time curves for steels A, B, C and D are plotted in Fig. 1 (where Pn = ( P , - Pequi)/Pq- Pequi); tot is the resistivity after aging time t, pq the resistivity of specimens quenched from solution temperature a n d Pequi the resistivity after a 1 h aging time). It is evident from these curves that, while resistivity on isothermal aging initially drops rapidly with little change thereafter, alloying additions had a significant influence on the rate of change in resistivity.
i~ 0.6
~ 0.~
O. o
=
o.2
Q
b
J
b
o 2
3 Logf lime- seconds
Fig. 1. Plot of normalized resistivity vs. aging time for steels A, B, C and D (isothermal aging temperature, 650 °C): a, 100% ferrite; b, onset of pearlite formation.
TABLE 2
Austenite grain sizes of steels
Specimen
A
B
C
D
Meanaustenite grain size (/zm)
293_+44
297_+46
218_+33
314_+47
The addition of silicon (alloy C) speeds the drop in resistivity, whereas the addition of manganese (alloy B) reduces the decrease. This effect is a direct indication of how these elements modify the transformation behaviour of steels and is clearly evident from the time each steel requires to reach the end of the 7 - a transformation, as shown in Fig. 1. Additionally, the smaller prior austenite grain size of steel C compared with those in steels A, B and D, given in Table 2, is also expected to make a contribution to the faster rate of transformation. Both the end of the ~ - a transformation and the prior austenite grain size in these steels were determined by optical metallography. A correlation between the resistivity and volume percentage of martensite for steels B and D is shown in Figs. 2 and 3. A strong linear correlation is observed and the slopes in each case are similar, 0.2 and 0.18 for steels B and D respectively. Similar plots for steels A and C could not be drawn as the 7 - a transformation in these alloys at 650 °C is very rapid.
191 90
E
I
I
I
I
30
8o
-b u~
~ 70 20
60
I
20
I 40
t~. t-o~
I
I
60
BO
100
volume°/, mtlrtensite
Fig. 2. Plot of resistivity v s . volume percentage of martensite for steel B, isothermallyaged at 650 °C.
I
I
I
10 *A IB -C *O
I
0 t-~/'-~
160
1
I
I
2
3 Log t lime-
seconds
Fig. 4. Plot of change in resistivity v s . aging time for steels A, B, C and D: a, 100%ferfite: b, onset of pearlite formation.
~5o > Lu_ .to
140 136
I
I
I
I
20
/*0
60
B0
votume%
100
mt~r'tensite
Fig. 3. Plot of resistivity v s . volume percentage of martensite for steel D, isothermallyaged at 650 °C. T h e curves for the change in resistivity with the isothermal aging time for steels A, B, C and D are plotted in Fig. 4. T h e change Ap in resistivity is calculated from A p = pq - Pt
It is clear from this set of curves, that the change in resistivity in all the alloys is approximately similar and lies between 26 and 30 nf2 m. A n examination of the extraction replicas indicated that, in all the alloys, vanadium carbide precipitation coincides with the start of the austenite transformation and occurred only in the ferrite. From the examination of thin foils, it was observed that precipitation formed in interphase (Fig. 5), random and fibrous precipitation (Fig. 6) morphologies. It has not been possible to establish when vanadium carbide precipitation reaches comple-
Fig. 5. Transmission electron micrograph from steel B isothermally aged for 9.5 rain at 650 °C, showing interphase carbide precipitation. tion. However, it is evident from optical metallography that pearlite formed in these alloys. Optical metallography (Fig. 7) also shows that pearlite formation is not concurrent with completion of the 7 - a transformation.
192
with omissions for the contributions from manganese plus silicon, from manganese and from silicon respectively. Under the present experimental conditions, neither manganese nor silicon forms carbides, and the change in resistivity due to vanadium carbide precipitation for all the alloys studied may therefore be written as
Ap = (apv + bpc)vc + Npdisln
Fig. 6. Transmission electron micrograph from steel B isothermally aged for 9.5 min at 650 °C, showing fibrous carbide precipitation.
4. Discussion
4.1. Contribution to resistivity The resistivity Pa~loyof dilute alloys containing one or more alloying elements may be expressed using Matthiessen's rule as Palloy = Ppure +
aPA + bloB
( 1)
w h e r e Ppure is the resistivity contribution from the
pure solvent metal, and aPA and bpB are resistivity contributions from solute atoms A and B of concentrations a and b respectively. By combining this rule with the hypothesis that, on isothermally aging a solid solution, the decrease Ap in resistivity is due to precipitation of solute atoms, removal of dislocations and grain growth, Ap may be expressed as Ap = (apA+ bpn)ppt + Npd,s,, + Pgb
(2)
where (apA+ bpB)ppt is the combined resistivity contribution from A and B to form a precipitate AB, a and b are concentrations of atoms A and B respectively, N is the number of dislocations, Pdis~n is the contribution from dislocations and Pgb is the contribution from changes in grain size. In the present study, for solution-treated-andquenched alloy D, eqn. ( 1 ) may be written as Palloy "~-PFc + PC + PV + PMn + PSi + Pdisln
(3)
For solution-treated-and-quenched alloys A, B and C, similar expressions may be written but
(4)
where (aPv + bpc)vc is the resistivity contribution from a wt.% V and b wt.% C to form stoichiometric VC- or VgC3-type precipitates; in the case of VC precipitation, b = 0 . 2 4 a and, for V4C3, b=0.18a. However, as there is more carbon than is required for either VC- or V4C3-type precipitates, the excess carbon may combine with iron to form Fe3C, and also some may remain in solution. At the level of manganese present in both steel B and steel D, some partitioning of manganese between the cementite and ferrite phases can also be expected. However, from the data given by Razik et al. [9] on their work on the partitioning of manganese in eutectoid steels, in which they found that the partitioning temperatures for 1.08 wt.% Mn and 1.8 wt.% Mn are 683 °C and 645 °C respectively, it is evident that under the present conditions of 1.5 wt.% Mn and 650 °C, no partitioning will occur. Thus the total resistivity change at equilibrium may be written as
Peq,i=(apv+ bpc)vc +(Cpc)F¢c+(epc)ci.sol.
(5)
where b + c + e is the total weight percentage of carbon in the alloy. To predict the rates of transformation and of precipitation using eqn. (5), accurate information about the resistivity contribution from components such as vanadium, carbon and dislocations in iron at 77 K is required.
4.2. Resistivity contribution from the iron lattice and carbon-free martensite The average resistivity contribution from wellannealed iron at 77 K has been given by both Speich [10] and by Fujita and Damask [11] as 8 nQ m. This value includes contributions from grain boundaries as well as from residual substitutional impurities. By extrapolating the electrical resistivity data for low carbon martensites to zero carbon, Speich [10] derived a contribution of 4 ng~ m for carbonfree martensite.
193
Fig. 7. Optical micrographs showing the microstructure of samples of steel B isothermally aged at 650 °C for (a) 1.5 min, (b) 4.5 min, (c) 9.5 min and (d) 59.5 min.
4.3. Resistivity contribution from vanadium T h e resistivity data reported in the literature for vanadium in iron at 77 K are given in Table 3. T h e contributions in both references are in good agreement, and a value of 27 nQ m is adopted for the present case.
4.4. Resistivity contribution from carbon in iron Data for the resistivity contribution from carbon in iron are given in Table 4. It is clear from these data that the contribution from carbon is dependent both on the phase and
on the carbon content. In martensites containing 0.2 wt.% C and less, the contribution is about a third that in martensites with more than 0.2 wt.% C. In ferrite, the contribution is independent of carbon content but is much higher than for low carbon martensites and almost the same as in high carbon martensite. Speich [10] has shown that this difference arises from segregation of carbon to dislocations or lath boundaries in low carbon martensites, whereas little segregation occurs in martensites containing more than 0.2 wt.% C. A similar inference had earlier been
194 TABLE 3 at 77 K
Resistivity contribution from vanadium in iron
TABLE 5
Comparison of observed and calculated resis-
tivities
Contribution a (nf2 m (wt.%)-')
Reference
Steel
Pobs (nf2 m)
Pcalc (nQ m)
26.8 27.0
[12] [13]
A B C D
29.9 26.6 29.8 27.8
30.4 30.4 28.4 28.2
~Calculated by interpolating between resistivity of pure iron taken as 8 nf~ m and the reported F e - V data.
TABLE 4
Resistivity contribution data from carbon in iron at 77 K (data obtained at 23 °C)
( bntribution (nQ m (wt.% C)
1)
Quenching temperature
Reference
(°c) <0.2 wt.% C
>0.2 wt.% C
100 -278 228 305 200
295 218.7 --
1000 1100 --
[1 O] [14] [15]
--
--
[16]
---
724 712
[17] [11]
Takamura et al. [18] confirmed that Matthiessen's rule holds for Fe-C alloys in the range 7 7 - 3 0 0 K.
drawn by King and Glover [14] for low carbon martensites. However, there is wide variation in the data for carbon contributions in ferrite, ranging from as low as 200 up to 305 n ~ m (wt.% C)-1. This variation is difficult to explain, although errors from a number of sources can arise, e.g. the accuracy of techniques used to measure specimen sizes and voltage drops, variations in temperature, impurity levels and heat treatment, adding up to a considerable total error. By applying Matthiessen's rule and using the carbon contribution values reported by King and Glover [14], Swartz and Cuddy [16] and Fujita and Damask [11], the resistivities of a 1.33 wt.% C martensite of 303 nff2 m, 315 nf2 m and 278 nf2 m respectively are predicted. (These values take into account a contribution of 4 n ~ m from martensite and dislocations, together with a lattice contribution of 8 nff2 m.) The predicted values are in reasonable agreement with the mean experimental value of 325 nff2 m for such a steel, as reported by Gavin [19]. Despite the many sources of error, it appears that the resistivity contributions from carbon in ferrite found by Speieh [10], Arajs [15] and Waganblast and Arajs [17] are somewhat high.
4.5. Evaluation of resistivity Using the appropriate functions in eqn. (3), for steel A, and contributions from vanadium of 27 nQ m (wt.%)-1, from iron of 8 nQ m (wt.%)-1, from carbon of 100 nQ m (wt.%) -l and from martensite or dislocations 4 nf2 m, the resistivity Pa of the solution-treated-and-quenched alloy is predicted to have a value of 38.4 nff2 m. This value is lower than the value of 44.4 ( + 2%) nf2 m determined experimentally. (Similar comparisons of calculated resistivity values for steel B, C and D have not been made because the contributions due to silicon and manganese are unavailable.) The discrepancy between the calculated and experimentally determined resistivities for alloy A may be explained by the omission of any impurity contributions in eqn. (3) and also by deviations from Matthiessen's rule, which are known to occur. The contributions from impurities and also from alloying additions not contributing to the precipitation reaction may be eliminated by considering only the changes in resistivity after the alloys had been aged. A comparison between the calculated changes in resistivity (eqn. (5)) and the observed changes pq - - P e q u i in resistivity is given in Table 5. With the exception of alloy B, an agreement within experimental error of about _+4% is found. This confirms that the observed changes in resistivity data can be interpreted using Matthiessen's rule, together with the hypothesis discussed above to follow carbide precipitation. Table 6 presents the calculated contributions PcaJcfrom VC precipitation (first term in eqn. (4)), the estimated contributions Pest from VC precipitation (calculated by subtracting from the total observed resistivity Pobs, the contribution P'c from excess carbon and the contribution Pdisln from martensite, i.e. Pest "= Pnbs -- P t C --/Odisln)" Also given in this table is the percentage of VC precipitation, calculated by taking the ratio of the estimated to calculated contributions from VC. As there is an uncertainty about the composition of
195
TABLE 6 Steel
A B C D
Calculated and estimated resistivity data for vanadium carbides ~ (. ) precipitate
VC precipitate P ~al,(nff2 m)
P ~t (nf2 m)
Amount (%)
p ~,~ (n• m)
p~ (nf2 ml
Amount (%)
23.3 23.3 23.3 22.8
22.7 19.3 24.6 22.4
97 83 106 98
20,6 20.6 20.6 20.1
2(I.0 16.7 21.9 19.8
97 81 106 99
vanadium carbide, a similarly calculated set of values for V4C 3 is given in Table 6. It is evident from this table that either type of vanadium carbide precipitation can reach completion after aging for about 1 h at 650 °C, except in the case of steel B. This indicates that, while vanadium carbide precipitation is observed to start with 7 - a transformation, precipitation does not reach completion until well after the completion of the y - a transformation. In Table 7 are given the precipitation percentage values considering both VC- and V4C3-type carbides, calculated by taking the ratio of change in resistivity at the finish of the y - a transformation (also given in Table 7) to the calculated contribution from the carbide precipitates (Table 6). The change in resistivity at y - a transformation is estimated by readings from Fig. 4, and subtracting from these values the contribution from martensite of 4 nQ m. It is clear that from the set of values given in Table 7, the bulk of carbide precipitation occurs concurrently with y - a transformation. However, a higher volume fraction of V4C~ precipitation is predicted for steels A, B and D while, for steel C, completion of precipitation should result. The precipitation percentage for the VC composition is lower and is inferred to be the predominant carbide, since the onset of pearlite observed from optical examination (Fig. 4) is not concurrent with the completion of the 7 - a transformation. This observation can be explained by the possibility that vanadium and carbon are present at the 7 - a interface as clusters which suppress pearlite nucleation. The formation of pearlite requires carbon to dissociate first from clusters [21 ]. The probability that the VC type of carbide predominates can also be inferred by considering the excess carbon available in the steels and the observed pearlite volume percentage in the aged alloys. Calculated values of excess carbon after either VC or V4C3 precipitation together with the mean pearlite percentage determined by optical
TABLE 7 Calculated percentage precipitation for vanadium carbides Steel
Estimated change in p at 7- a finish due to precipitate
A m o u n t (%) of precipitate
(nt2 m) A B C D
18.0 16.5 21.0 16.0
77 71 90 70
87 8{} I {}0 80
metallography are given in Table 8. From the determined pearlite percentages, and on the assumption that the eutectoid concentration of 0.8 wt.% C in the Fe-C system is unaffected by the alloying additions made, the carbon combined into pearlite can be calculated. From these data, the amount of carbon in solution can be determined. Both these sets of values are also given in Table 8, and it is evident that, in the case of VC3 precipitation, the carbon in solution is much higher than can be expected at 650 °C but quite reasonable for the case of VC precipitation. The pearlite fraction is observed to be highest in steel B (Table 8) with the calculated percentage of vanadium carbide precipitation of only about 80%. In addition, the calculated vanadium carbide precipitation is a maximum in steels C and D, which also show a minimum pearlite formation. These observations are consequences of the alloying additions; silicon stabilizes ferrite, speeding up the y - a transformation and thereby allowing less time for vanadium carbide formation but more time for pearlite formation. Manganese, however, by stabilizing austenite, considerably slows down the 7 - a transformation, allowing more time for vanadium carbide formation and consequently developing a lower pearlite volume fraction, in addition to allowing the retention of more carbon in solution. From the above discussion, it can be concluded that the point of completion of alloy carbide
196 TABLE 8
Steel
A B C D
Calculated percentages of excess carbon and carbon combined as pearlite
Amount (wt.% C) of excess C
Mean amount of pearlite (experimental)
vc
v4c,
(%)
0.03 0.03 0.01 0.01
0.06 0.06 0.04 0.04
1.8 4.3 1.0 1.0
precipitation cannot be deduced directly from electrical resistivity measurements, since t h e reduction in resistivity in the final stages of aging is due to concurrent alloy and iron carbide precipitation, and there is also an uncertainty in the composition of alloy carbide [20].
Amount (wt.% C) of C combined in pearlite
Amount (wt.% C) of C in solution
0.014 0.034 0.008 0.008
VC
v~C~
0.016 -0.002 0.002
0.045 0.025 0.031 0.032
500
400 o :> -r
Q;
o
E
*
•
÷
° •
4.6. Microhardness values
The inference from electrical resistivity data considerations that, in the present experiments, VC predominates and that the completion of vanadium carbide precipitation takes place well after the end of the y - a transformation suggests that an accompanying increase in the hardness of the ferrite phase must occur. Peak hardness should coincide with maximum precipitation volume fraction. This is indeed the case shown in Fig. 8, in which are plotted the microhardness values as a function of aging times for both the ferrite and the martensite phases in steels. It is clear from this figure that the maximum hardness of the ferrite phase in steels A and B occurs after aging for 20 min, whereas the 7 - a transformation was complete after aging for about 2 min and 10 min respectively. However, for steels C and D, the maximum hardness is observed after aging for about 5 min although, in the case of steel C, the 7 - a transformation finished after aging for between 0 and 0.5 rain, and in steel D this completion occurred after aging for about 4 min. In addition to affirming the information ascertained from electrical resistivity data, the time required for peak hardness of the ferrite or the maximum volume fraction of alloy carbide precipitation provides further evidence of the influence of silicon and manganese on precipitation kinetics. With respect to steel A, which contains neither of these additions, the absence of manganese appears to have little effect on carbide precipitation kinetics. On the contrary, the silicon addition alone increases the rate by about four times in steel C. However, this estimated increased
300
Q
!
b 2
°A ,B "C "D 200
I
I
2 time
log t
3
- seconds
Fig. 8. Plot of microhardness vs. aging time of both the martensite and the ferrite phases occurring in isothermally aged steels A, B, C and D.
rate may be somewhat low as there must be added a contribution due to the increased rate of transformation through the smaller prior austenite grain size in steel C (Table 2). The addition of manganese and silicon together produces a similar increase in the precipitation rate to that of silicon on its own, i.e. silicon not only speeds up the y - a transformation but also increases the precipitation rate, while manganese only slows down the y - a kinetics and has no effect on precipitation kinetics. Moreover, when manganese is added together with silicon, the effect of silicon predominates, both in accelerating the 7 - a transformation and also in increasing the precipitation rate of vanadium carbide, although the prior austenite grain size appears to be unaffected by the silicon addition (Table 2). The increase in the hardness of the martensite found in steels B and D can be explained by assuming a partitioning of carbon between austenite and ferrite during aging. This phenomenon has also been reported by Sakuma and Honeycombe [22] during isothermal aging of
197
Fe-Nb-C alloys, where a considerable hardening of the martensite was observed. In the present alloys, the hardening is not so pronounced as only a small fraction of the carbon remains free to partition after carbide precipitation which, as indicated above, occurs predominantly and concurrently with austenite transformation. The fraction of carbon free to partition is shown in Table 8, i.e. about 20% in the case of steel B and about 10% for steel D. The slightly higher carbon composition of steel B, combined with the reduced austenite transformation rate arising from the addition of manganese, also provides an explanation for the higher hardness of the martensite recorded in this steel. On the contrary, in the study of Sakuma and Honeycombe [22], from chemical compositional considerations it can be shown that only a small fraction (about 6 wt.%) of the carbon precipitates out as NbC, leaving a considerable amount of carbon to partition freely between austenite and ferrite, thereby markedly increasing the hardness of the martensite formed during isothermal aging. 4. 7. Particle size distributions Changes in electrical resistivity are insensitive to particle size and therefore no information regarding this parameter, or the changes in particle size during aging, can be deduced from resistivity measurements, to explain the softening of the ferrite observed on continued aging, which
arises from particle coarsening at constant volume fraction. The typical particle size distributions observed in the present alloys, obtained using extraction replicas, are illustrated in the histograms shown in Figs. 9-11. These histograms show percentage particle frequency against size for steel D aged for 1.5, 9.5 and 59.5 min. For an isothermal aging time of about 2 min, the precipitation appears homogeneous with a unimodal particle size (Fig. 9) and a mean particle size distribution of about 2.0 nm. After aging for /.0
~ 30 z LLI 0 L.iJ LL
2O
t.,u u (_1
0
I I 15 2.5 3.5 SIZE (nm}
05
(a) I
l
I
4.5
I
I
I
I
I
I
I
20
50 z bJ
g
z LtJ
k~
0 t.u
U_ W (D
~25
~D
n
oV 0.5
1.5
SIZE
25
\ F 0.8
3.5
/..5
(nm)
Fig. 9. Histogram of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 1.5 min at 650 °C.
2.0
I 32
I I I 4.; 5.6 6.B SIZE (nm)
B.O
9.2 10./.
(b) Fig. 10. Histograms of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 9.5 min at 6 5 0 °C.
198 40
I
I
distribution of particles can be markedly affected during the preparation of extraction replicas. A similar inhomogeneous particle size distribution is also observed after aging for about 60 min (Fig. 11). In addition, the frequency of the coarser particles (Fig. 1 l(b)) has also increased. Since it was not possible to estimate the particle volume fraction from extraction replicas [23], from the variation in the hardness of ferrite observed (Fig. 8) it can be inferred that the finer particles in the early stages of aging arise owing to nucleation with little growth and that an increasing particle volume fraction results in an increase in the hardness. In the later part of aging, however, a decrease in hardness occurs because of particle growth at constant particle volume fraction. The existence of finer particles at this stage can be explained by the particle dissolution and the growth of the more stable larger particles.
I
j30 z LU 0 UJ
w 20 ...J (.3
< fl_
10
0
(a) 20
L
fl
I I 0.5 1.5 2.5 3.5 4.5 SIZE (nm)
I
I
I
I
I
I
I
I
\
5. Conclusions
Z IJA (21 W LI_
m 10 J
a_
oJ 0.8 2.0
I 3.2
I 4.4
I 5.4
I 6.B
I I ~ 1 8.0 9.2 10.4 11.6
S I Z E (nm)
(b) Fig. 11. Histograms of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 59.5 min at 650 °C.
about 10 min, the particle size distribution is observed to be inhomogeneous, with finer particles (Fig. 10(a)) mainly occurring within grains and coarser particles (Fig. 10(b)) at grain boundary regions. Particles near grain boundaries are coarser since early precipitation occurs in these regions, as the first transformation is initiated here. Hence these precipitates have an appreciably longer period for growth than those particles which nucleate at a later stage during aging. However, it must also be pointed out that the true
Alloy carbide precipitation during the isothermal aging of F e - C - V alloys at 650 °C starts simultaneously with austenlte transformation and occurs only in the ferrite phase. While the bulk of the carbide precipitation forms during the austenite-to-ferrite transformation, a significant proportion takes place after the finish of transformation. The kinetics of precipitation are strongly affected by alloying additions through their respective effect on the time-temperaturetransformation curves. Silicon speeds up the transformation, with the consequence that most vanadium carbide precipitation occurs within an aging time of 2 min. Silicon also increases the precipitation rate such that complete precipitation occurs after a further aging of about 3 min. On the contrary, manganese, by slowing down the austenite transformation, allows the bulk of the precipitation to occur after aging for 10 rain and complete precipitation after a further aging of about 10 rain. Moreover, other than affecting the austenite transformation behaviour, the addition of manganese has little influence on precipitation kinetics. The addition of manganese and silicon together results in an intermediate transformation rate such that bulk of the precipitation occurs after aging for about 7 min. The predominating effect of silicon alone in increasing the rate of
199
precipitation by about four times is little affected in steel D with combined addition. Acknowledgments The authors are grateful to the Science and Engineering Research Council for funding this work and to the Head of the Department of Metallurgy, University of Strathclyde, for providing research facilities. References 1 W. B. Morrison, J. Iron Steel Inst., London, 102 (1963) 317. 2 R. Phillips and J. A. Chapman, J. Iron Steel Inst., London, 204(1966) 615. 3 K. Bunghardt, K. Kind and W. Oelsen, Arch. Eisenhiittenwes., 26 (1956) 61. 4 R. K. Amin, M. Korchynsky and E B. Picketing, Met. Technol., 8 ( 1981 ) 250. 5 M. G. Akben, 1. Weiss and J. J. Jonas, Acta Metall., 29 (1981) 111. 6 T. N. Baker, Heat Treatment '73, Metals Society, London, 1975, p. 13. 7 W.B. Morrison, B. Mintz and R. C. Cochrane, Controlled Processing of High Strength Low Alloy Steels, Preprints BSC Conf., 1976.
8 W.G. Hall and T. N. Baker, Met. Sci., 15 ( 1981 ) 447. 9 N. A. Razik, G. W. Lorimer and N. Ridley, Acta Metall., 22 (1974) 1249. 10 G. R. Speich, Trans. Metall. Soc. AIME, 245 (1969) 2553. 11 F. E. Fujita and A. C. Damask, Acta Metall., 12 (1964) 331. 12 A. V. Cheremuslkina and M. I. Koraleva, Soy. Phys.-Solid State, 5 (1963) 330. 13 M. M. Yang and A. D. Krawitz, Metall. Trans. A, 15 (1984) 1545. 14 H. W. King and S. G. Glover, J. Iron Steel lnst., London, 188(1958)61. 15 S. Arajs, Phys. Status Solidi, 33 (1969) 731. 16 J. C. Swartz and L. J. Cuddy, J. Phys. Chem. Solids, 32 (1971) 385. 17 H. Waganblast and S. Arajs, J. Appl. Phys., 3 (1968) 5. 18 J. Takamura, I. Takahashi and M. Amano, Trans. Iron Steel Inst. Jpn., 9 (1969) 216. 19 P.J. Gavin, Scr. Metall., 3 (1969) 741. 20 W. Roberts and A. Sandberg, Report IM 1489, 1980 (Swedish Institute for Metals Research). 21 R. C. Sharma and G. R. Purdy, Metall. Trans., 4 (1973) 23O4. 22 T. Sakuma and R. W. K. Honeycombe, Met. Sci., 18 (1984) 449. 23 T. N. Baker and H. S. Ubhi, in G. W. Lorimer (ed.), Analytical Electron Microscopy, Proc. Electron Microscopy and Analysis Group, Institute of Metals, London, 1988, p. 135.
189
The Influence of Manganese and Silicon on the Precipitation of Vanadium Carbide in Steel H. S. UBHI* and T. N. B A K E R
Division of Metallurgy and Engineering Materials, Universityof Strathclyde, Glasgow, G I 1XN (U.K.) (Received November 2, 1987; in revised form October 31, 1988)
Abstract
Four vacuum-melted steels with a nominal composition of Fe-O.12wt. %C-0.46wt. %V contained additions of 1.5 wt. % Mn (steel B), 0.3 wt. % Si (steel C) and 1.5 wt. % Mn and 0.3 wt. % Si (steel D). Steel A, with no manganese or silicon additions, was the control steel. After hot rolling, some of the material was drawn to wire, prior to solution treating at 1250 °C followed by aging for different times at 650 °C. Resistivity, hardness and metallography were used to follow the precipitation kinetics of vanadium carbide. Alloy carbide precipitation at 650 °C was found to start simultaneously with austenite transformation and occurred only in the ferrite phase. Silicon increased the rate of precipitation by approximately four times, but manganese, other than affecting the rate of austenite transformation, had little influence on the alloy precipitation kinetics. The combined addition of manganese and silicon resulted in an intermediate rate of transformation which influenced the precipitation kinetics of vanadium carbide.
tion strengthening, further tempering may be required [6, 7]. In order, therefore, to optimize the benefit from precipitation strengthening, a knowledge of the precipitation kinetics is of vital importance. However, few data on this topic are available; hence the object of this work is to investigate the kinetics of vanadium carbide precipitation in 0.12wt.%C-0.46wt.%V steels as well as the effects of manganese and silicon additions when added singly and together. This has been done by exploiting the extreme sensitivity of electrical resistivity changes in metals due to processes such as precipitation and dislocation annealing. In addition, transmission electron microscopy has been used to give structural information as a complement to the electrical resistivity data. This study is confined to precipitation during isothermal aging and is an extension of previous studies on the base steel [8]. Vacuum-melted microalloyed steels have been used to study the influence of manganese and silicon on the rate of isothermal precipitation of vanadium carbide in the ferrite phase. The four steels (Table 1) used had a nominal composition
I. Introduction
TABLE 1
High strengths in low alloy steels are achieved primarily by a combination of grain refinement and precipitation hardening [1]. Whereas the former can be produced by using stringent rolling procedures [2], the latter usually occurs both during and after rolling [3-5]. However, it is evident that, to reach the full potential of precipita-
Steel
*Present address: Department of Mechanical Engineering, University of Nairobi, Kenya. 0921-5093/89/$3.50
A B C D
Alloy compositions
Amount (wt.%) offollowing elements V
C
Si
Mn
0.46 0.46 0.46 0.45
0.14 0.14 0.12 0.12
0.01 0.01 0.28 0.30
0.01 1.5 0.01 1.5
The steels also contained 0.004 wt.% P, 0.003 wt.% S, less than 0.01 wt.% Cr, less than 0.005 wt.% Mo, less than 0.01 wt.% Ni, less than 0.005 wt.% AI, less than 0.001 wt.% N and less than 0.006 wt.% O, with the balance iron.
© Elsevier Sequoia/Printed in The Netherlands
190
of Fe-0.12wt.%C-0.46wt.%V. Manganese at the 1.5 wt.% and silicon at the 0.3 wt.% levels were present singly or in combination. The carbon-totransition-metal ratio was chosen to be approximately equivalent to the stoichiometric VC ratio.
1.0 ~
'
~
oA OB
oC +D 0.8
>,
2. Experimental details The vacuum-melted steels, the compositions of which are given in Table 1, were supplied in the form of hot-rolled 32 mm 2 square bars. The bars were cut into 150 mm lengths, hot rolled to rod of 10 mm diameter and then cold drawn to wire of 3 mm diameter. Specimens of 100 mm length were cut from this wire and preheat treated in encapsulated silica tubes for 5 h at 1200 °C followed by air cooling. They were then re-encapsulated, solution treated for 15 min at 1250 °C followed by isothermal aging at 650 °C in a tin bath for specified aging times. After this time lapse, the specimens were quickly lifted out of the tin bath, the silica tube was broken open and the specimen was allowed to drop into a quenching tank containing a medium of 4% aqueous brine. This operation was completed within 15 s. The voltage drop across the aged specimens was measured in liquid nitrogen, using the classical four-point method. A current of 0.5 A from a Hewlett-Packard G177B source was passed through the specimen in series with a 0.20 f2 standard resistance held in a constant-temperature bath. The potential drop across both the specimen and the standard resistance was measured with a Solartron A210 digital voltmeter. Resistivity values were obtained to an accuracy of + 2%. Both thin foil and carbon extraction replicas were prepared by standard techniques and examined using a Philips EM400T transmission electron microscope. 3. Results Normalized resistivity p , vs. isothermal aging time curves for steels A, B, C and D are plotted in Fig. 1 (where Pn = ( P , - Pequi)/Pq- Pequi); tot is the resistivity after aging time t, pq the resistivity of specimens quenched from solution temperature a n d Pequi the resistivity after a 1 h aging time). It is evident from these curves that, while resistivity on isothermal aging initially drops rapidly with little change thereafter, alloying additions had a significant influence on the rate of change in resistivity.
i~ 0.6
~ 0.~
O. o
=
o.2
Q
b
J
b
o 2
3 Logf lime- seconds
Fig. 1. Plot of normalized resistivity vs. aging time for steels A, B, C and D (isothermal aging temperature, 650 °C): a, 100% ferrite; b, onset of pearlite formation.
TABLE 2
Austenite grain sizes of steels
Specimen
A
B
C
D
Meanaustenite grain size (/zm)
293_+44
297_+46
218_+33
314_+47
The addition of silicon (alloy C) speeds the drop in resistivity, whereas the addition of manganese (alloy B) reduces the decrease. This effect is a direct indication of how these elements modify the transformation behaviour of steels and is clearly evident from the time each steel requires to reach the end of the 7 - a transformation, as shown in Fig. 1. Additionally, the smaller prior austenite grain size of steel C compared with those in steels A, B and D, given in Table 2, is also expected to make a contribution to the faster rate of transformation. Both the end of the ~ - a transformation and the prior austenite grain size in these steels were determined by optical metallography. A correlation between the resistivity and volume percentage of martensite for steels B and D is shown in Figs. 2 and 3. A strong linear correlation is observed and the slopes in each case are similar, 0.2 and 0.18 for steels B and D respectively. Similar plots for steels A and C could not be drawn as the 7 - a transformation in these alloys at 650 °C is very rapid.
191 90
E
I
I
I
I
30
8o
-b u~
~ 70 20
60
I
20
I 40
t~. t-o~
I
I
60
BO
100
volume°/, mtlrtensite
Fig. 2. Plot of resistivity v s . volume percentage of martensite for steel B, isothermallyaged at 650 °C.
I
I
I
10 *A IB -C *O
I
0 t-~/'-~
160
1
I
I
2
3 Log t lime-
seconds
Fig. 4. Plot of change in resistivity v s . aging time for steels A, B, C and D: a, 100%ferfite: b, onset of pearlite formation.
~5o > Lu_ .to
140 136
I
I
I
I
20
/*0
60
B0
votume%
100
mt~r'tensite
Fig. 3. Plot of resistivity v s . volume percentage of martensite for steel D, isothermallyaged at 650 °C. T h e curves for the change in resistivity with the isothermal aging time for steels A, B, C and D are plotted in Fig. 4. T h e change Ap in resistivity is calculated from A p = pq - Pt
It is clear from this set of curves, that the change in resistivity in all the alloys is approximately similar and lies between 26 and 30 nf2 m. A n examination of the extraction replicas indicated that, in all the alloys, vanadium carbide precipitation coincides with the start of the austenite transformation and occurred only in the ferrite. From the examination of thin foils, it was observed that precipitation formed in interphase (Fig. 5), random and fibrous precipitation (Fig. 6) morphologies. It has not been possible to establish when vanadium carbide precipitation reaches comple-
Fig. 5. Transmission electron micrograph from steel B isothermally aged for 9.5 rain at 650 °C, showing interphase carbide precipitation. tion. However, it is evident from optical metallography that pearlite formed in these alloys. Optical metallography (Fig. 7) also shows that pearlite formation is not concurrent with completion of the 7 - a transformation.
192
with omissions for the contributions from manganese plus silicon, from manganese and from silicon respectively. Under the present experimental conditions, neither manganese nor silicon forms carbides, and the change in resistivity due to vanadium carbide precipitation for all the alloys studied may therefore be written as
Ap = (apv + bpc)vc + Npdisln
Fig. 6. Transmission electron micrograph from steel B isothermally aged for 9.5 min at 650 °C, showing fibrous carbide precipitation.
4. Discussion
4.1. Contribution to resistivity The resistivity Pa~loyof dilute alloys containing one or more alloying elements may be expressed using Matthiessen's rule as Palloy = Ppure +
aPA + bloB
( 1)
w h e r e Ppure is the resistivity contribution from the
pure solvent metal, and aPA and bpB are resistivity contributions from solute atoms A and B of concentrations a and b respectively. By combining this rule with the hypothesis that, on isothermally aging a solid solution, the decrease Ap in resistivity is due to precipitation of solute atoms, removal of dislocations and grain growth, Ap may be expressed as Ap = (apA+ bpn)ppt + Npd,s,, + Pgb
(2)
where (apA+ bpB)ppt is the combined resistivity contribution from A and B to form a precipitate AB, a and b are concentrations of atoms A and B respectively, N is the number of dislocations, Pdis~n is the contribution from dislocations and Pgb is the contribution from changes in grain size. In the present study, for solution-treated-andquenched alloy D, eqn. ( 1 ) may be written as Palloy "~-PFc + PC + PV + PMn + PSi + Pdisln
(3)
For solution-treated-and-quenched alloys A, B and C, similar expressions may be written but
(4)
where (aPv + bpc)vc is the resistivity contribution from a wt.% V and b wt.% C to form stoichiometric VC- or VgC3-type precipitates; in the case of VC precipitation, b = 0 . 2 4 a and, for V4C3, b=0.18a. However, as there is more carbon than is required for either VC- or V4C3-type precipitates, the excess carbon may combine with iron to form Fe3C, and also some may remain in solution. At the level of manganese present in both steel B and steel D, some partitioning of manganese between the cementite and ferrite phases can also be expected. However, from the data given by Razik et al. [9] on their work on the partitioning of manganese in eutectoid steels, in which they found that the partitioning temperatures for 1.08 wt.% Mn and 1.8 wt.% Mn are 683 °C and 645 °C respectively, it is evident that under the present conditions of 1.5 wt.% Mn and 650 °C, no partitioning will occur. Thus the total resistivity change at equilibrium may be written as
Peq,i=(apv+ bpc)vc +(Cpc)F¢c+(epc)ci.sol.
(5)
where b + c + e is the total weight percentage of carbon in the alloy. To predict the rates of transformation and of precipitation using eqn. (5), accurate information about the resistivity contribution from components such as vanadium, carbon and dislocations in iron at 77 K is required.
4.2. Resistivity contribution from the iron lattice and carbon-free martensite The average resistivity contribution from wellannealed iron at 77 K has been given by both Speich [10] and by Fujita and Damask [11] as 8 nQ m. This value includes contributions from grain boundaries as well as from residual substitutional impurities. By extrapolating the electrical resistivity data for low carbon martensites to zero carbon, Speich [10] derived a contribution of 4 ng~ m for carbonfree martensite.
193
Fig. 7. Optical micrographs showing the microstructure of samples of steel B isothermally aged at 650 °C for (a) 1.5 min, (b) 4.5 min, (c) 9.5 min and (d) 59.5 min.
4.3. Resistivity contribution from vanadium T h e resistivity data reported in the literature for vanadium in iron at 77 K are given in Table 3. T h e contributions in both references are in good agreement, and a value of 27 nQ m is adopted for the present case.
4.4. Resistivity contribution from carbon in iron Data for the resistivity contribution from carbon in iron are given in Table 4. It is clear from these data that the contribution from carbon is dependent both on the phase and
on the carbon content. In martensites containing 0.2 wt.% C and less, the contribution is about a third that in martensites with more than 0.2 wt.% C. In ferrite, the contribution is independent of carbon content but is much higher than for low carbon martensites and almost the same as in high carbon martensite. Speich [10] has shown that this difference arises from segregation of carbon to dislocations or lath boundaries in low carbon martensites, whereas little segregation occurs in martensites containing more than 0.2 wt.% C. A similar inference had earlier been
194 TABLE 3 at 77 K
Resistivity contribution from vanadium in iron
TABLE 5
Comparison of observed and calculated resis-
tivities
Contribution a (nf2 m (wt.%)-')
Reference
Steel
Pobs (nf2 m)
Pcalc (nQ m)
26.8 27.0
[12] [13]
A B C D
29.9 26.6 29.8 27.8
30.4 30.4 28.4 28.2
~Calculated by interpolating between resistivity of pure iron taken as 8 nf~ m and the reported F e - V data.
TABLE 4
Resistivity contribution data from carbon in iron at 77 K (data obtained at 23 °C)
( bntribution (nQ m (wt.% C)
1)
Quenching temperature
Reference
(°c) <0.2 wt.% C
>0.2 wt.% C
100 -278 228 305 200
295 218.7 --
1000 1100 --
[1 O] [14] [15]
--
--
[16]
---
724 712
[17] [11]
Takamura et al. [18] confirmed that Matthiessen's rule holds for Fe-C alloys in the range 7 7 - 3 0 0 K.
drawn by King and Glover [14] for low carbon martensites. However, there is wide variation in the data for carbon contributions in ferrite, ranging from as low as 200 up to 305 n ~ m (wt.% C)-1. This variation is difficult to explain, although errors from a number of sources can arise, e.g. the accuracy of techniques used to measure specimen sizes and voltage drops, variations in temperature, impurity levels and heat treatment, adding up to a considerable total error. By applying Matthiessen's rule and using the carbon contribution values reported by King and Glover [14], Swartz and Cuddy [16] and Fujita and Damask [11], the resistivities of a 1.33 wt.% C martensite of 303 nff2 m, 315 nf2 m and 278 nf2 m respectively are predicted. (These values take into account a contribution of 4 n ~ m from martensite and dislocations, together with a lattice contribution of 8 nff2 m.) The predicted values are in reasonable agreement with the mean experimental value of 325 nff2 m for such a steel, as reported by Gavin [19]. Despite the many sources of error, it appears that the resistivity contributions from carbon in ferrite found by Speieh [10], Arajs [15] and Waganblast and Arajs [17] are somewhat high.
4.5. Evaluation of resistivity Using the appropriate functions in eqn. (3), for steel A, and contributions from vanadium of 27 nQ m (wt.%)-1, from iron of 8 nQ m (wt.%)-1, from carbon of 100 nQ m (wt.%) -l and from martensite or dislocations 4 nf2 m, the resistivity Pa of the solution-treated-and-quenched alloy is predicted to have a value of 38.4 nff2 m. This value is lower than the value of 44.4 ( + 2%) nf2 m determined experimentally. (Similar comparisons of calculated resistivity values for steel B, C and D have not been made because the contributions due to silicon and manganese are unavailable.) The discrepancy between the calculated and experimentally determined resistivities for alloy A may be explained by the omission of any impurity contributions in eqn. (3) and also by deviations from Matthiessen's rule, which are known to occur. The contributions from impurities and also from alloying additions not contributing to the precipitation reaction may be eliminated by considering only the changes in resistivity after the alloys had been aged. A comparison between the calculated changes in resistivity (eqn. (5)) and the observed changes pq - - P e q u i in resistivity is given in Table 5. With the exception of alloy B, an agreement within experimental error of about _+4% is found. This confirms that the observed changes in resistivity data can be interpreted using Matthiessen's rule, together with the hypothesis discussed above to follow carbide precipitation. Table 6 presents the calculated contributions PcaJcfrom VC precipitation (first term in eqn. (4)), the estimated contributions Pest from VC precipitation (calculated by subtracting from the total observed resistivity Pobs, the contribution P'c from excess carbon and the contribution Pdisln from martensite, i.e. Pest "= Pnbs -- P t C --/Odisln)" Also given in this table is the percentage of VC precipitation, calculated by taking the ratio of the estimated to calculated contributions from VC. As there is an uncertainty about the composition of
195
TABLE 6 Steel
A B C D
Calculated and estimated resistivity data for vanadium carbides ~ (. ) precipitate
VC precipitate P ~al,(nff2 m)
P ~t (nf2 m)
Amount (%)
p ~,~ (n• m)
p~ (nf2 ml
Amount (%)
23.3 23.3 23.3 22.8
22.7 19.3 24.6 22.4
97 83 106 98
20,6 20.6 20.6 20.1
2(I.0 16.7 21.9 19.8
97 81 106 99
vanadium carbide, a similarly calculated set of values for V4C 3 is given in Table 6. It is evident from this table that either type of vanadium carbide precipitation can reach completion after aging for about 1 h at 650 °C, except in the case of steel B. This indicates that, while vanadium carbide precipitation is observed to start with 7 - a transformation, precipitation does not reach completion until well after the completion of the y - a transformation. In Table 7 are given the precipitation percentage values considering both VC- and V4C3-type carbides, calculated by taking the ratio of change in resistivity at the finish of the y - a transformation (also given in Table 7) to the calculated contribution from the carbide precipitates (Table 6). The change in resistivity at y - a transformation is estimated by readings from Fig. 4, and subtracting from these values the contribution from martensite of 4 nQ m. It is clear that from the set of values given in Table 7, the bulk of carbide precipitation occurs concurrently with y - a transformation. However, a higher volume fraction of V4C~ precipitation is predicted for steels A, B and D while, for steel C, completion of precipitation should result. The precipitation percentage for the VC composition is lower and is inferred to be the predominant carbide, since the onset of pearlite observed from optical examination (Fig. 4) is not concurrent with the completion of the 7 - a transformation. This observation can be explained by the possibility that vanadium and carbon are present at the 7 - a interface as clusters which suppress pearlite nucleation. The formation of pearlite requires carbon to dissociate first from clusters [21 ]. The probability that the VC type of carbide predominates can also be inferred by considering the excess carbon available in the steels and the observed pearlite volume percentage in the aged alloys. Calculated values of excess carbon after either VC or V4C3 precipitation together with the mean pearlite percentage determined by optical
TABLE 7 Calculated percentage precipitation for vanadium carbides Steel
Estimated change in p at 7- a finish due to precipitate
A m o u n t (%) of precipitate
(nt2 m) A B C D
18.0 16.5 21.0 16.0
77 71 90 70
87 8{} I {}0 80
metallography are given in Table 8. From the determined pearlite percentages, and on the assumption that the eutectoid concentration of 0.8 wt.% C in the Fe-C system is unaffected by the alloying additions made, the carbon combined into pearlite can be calculated. From these data, the amount of carbon in solution can be determined. Both these sets of values are also given in Table 8, and it is evident that, in the case of VC3 precipitation, the carbon in solution is much higher than can be expected at 650 °C but quite reasonable for the case of VC precipitation. The pearlite fraction is observed to be highest in steel B (Table 8) with the calculated percentage of vanadium carbide precipitation of only about 80%. In addition, the calculated vanadium carbide precipitation is a maximum in steels C and D, which also show a minimum pearlite formation. These observations are consequences of the alloying additions; silicon stabilizes ferrite, speeding up the y - a transformation and thereby allowing less time for vanadium carbide formation but more time for pearlite formation. Manganese, however, by stabilizing austenite, considerably slows down the 7 - a transformation, allowing more time for vanadium carbide formation and consequently developing a lower pearlite volume fraction, in addition to allowing the retention of more carbon in solution. From the above discussion, it can be concluded that the point of completion of alloy carbide
196 TABLE 8
Steel
A B C D
Calculated percentages of excess carbon and carbon combined as pearlite
Amount (wt.% C) of excess C
Mean amount of pearlite (experimental)
vc
v4c,
(%)
0.03 0.03 0.01 0.01
0.06 0.06 0.04 0.04
1.8 4.3 1.0 1.0
precipitation cannot be deduced directly from electrical resistivity measurements, since t h e reduction in resistivity in the final stages of aging is due to concurrent alloy and iron carbide precipitation, and there is also an uncertainty in the composition of alloy carbide [20].
Amount (wt.% C) of C combined in pearlite
Amount (wt.% C) of C in solution
0.014 0.034 0.008 0.008
VC
v~C~
0.016 -0.002 0.002
0.045 0.025 0.031 0.032
500
400 o :> -r
Q;
o
E
*
•
÷
° •
4.6. Microhardness values
The inference from electrical resistivity data considerations that, in the present experiments, VC predominates and that the completion of vanadium carbide precipitation takes place well after the end of the y - a transformation suggests that an accompanying increase in the hardness of the ferrite phase must occur. Peak hardness should coincide with maximum precipitation volume fraction. This is indeed the case shown in Fig. 8, in which are plotted the microhardness values as a function of aging times for both the ferrite and the martensite phases in steels. It is clear from this figure that the maximum hardness of the ferrite phase in steels A and B occurs after aging for 20 min, whereas the 7 - a transformation was complete after aging for about 2 min and 10 min respectively. However, for steels C and D, the maximum hardness is observed after aging for about 5 min although, in the case of steel C, the 7 - a transformation finished after aging for between 0 and 0.5 rain, and in steel D this completion occurred after aging for about 4 min. In addition to affirming the information ascertained from electrical resistivity data, the time required for peak hardness of the ferrite or the maximum volume fraction of alloy carbide precipitation provides further evidence of the influence of silicon and manganese on precipitation kinetics. With respect to steel A, which contains neither of these additions, the absence of manganese appears to have little effect on carbide precipitation kinetics. On the contrary, the silicon addition alone increases the rate by about four times in steel C. However, this estimated increased
300
Q
!
b 2
°A ,B "C "D 200
I
I
2 time
log t
3
- seconds
Fig. 8. Plot of microhardness vs. aging time of both the martensite and the ferrite phases occurring in isothermally aged steels A, B, C and D.
rate may be somewhat low as there must be added a contribution due to the increased rate of transformation through the smaller prior austenite grain size in steel C (Table 2). The addition of manganese and silicon together produces a similar increase in the precipitation rate to that of silicon on its own, i.e. silicon not only speeds up the y - a transformation but also increases the precipitation rate, while manganese only slows down the y - a kinetics and has no effect on precipitation kinetics. Moreover, when manganese is added together with silicon, the effect of silicon predominates, both in accelerating the 7 - a transformation and also in increasing the precipitation rate of vanadium carbide, although the prior austenite grain size appears to be unaffected by the silicon addition (Table 2). The increase in the hardness of the martensite found in steels B and D can be explained by assuming a partitioning of carbon between austenite and ferrite during aging. This phenomenon has also been reported by Sakuma and Honeycombe [22] during isothermal aging of
197
Fe-Nb-C alloys, where a considerable hardening of the martensite was observed. In the present alloys, the hardening is not so pronounced as only a small fraction of the carbon remains free to partition after carbide precipitation which, as indicated above, occurs predominantly and concurrently with austenite transformation. The fraction of carbon free to partition is shown in Table 8, i.e. about 20% in the case of steel B and about 10% for steel D. The slightly higher carbon composition of steel B, combined with the reduced austenite transformation rate arising from the addition of manganese, also provides an explanation for the higher hardness of the martensite recorded in this steel. On the contrary, in the study of Sakuma and Honeycombe [22], from chemical compositional considerations it can be shown that only a small fraction (about 6 wt.%) of the carbon precipitates out as NbC, leaving a considerable amount of carbon to partition freely between austenite and ferrite, thereby markedly increasing the hardness of the martensite formed during isothermal aging. 4. 7. Particle size distributions Changes in electrical resistivity are insensitive to particle size and therefore no information regarding this parameter, or the changes in particle size during aging, can be deduced from resistivity measurements, to explain the softening of the ferrite observed on continued aging, which
arises from particle coarsening at constant volume fraction. The typical particle size distributions observed in the present alloys, obtained using extraction replicas, are illustrated in the histograms shown in Figs. 9-11. These histograms show percentage particle frequency against size for steel D aged for 1.5, 9.5 and 59.5 min. For an isothermal aging time of about 2 min, the precipitation appears homogeneous with a unimodal particle size (Fig. 9) and a mean particle size distribution of about 2.0 nm. After aging for /.0
~ 30 z LLI 0 L.iJ LL
2O
t.,u u (_1
0
I I 15 2.5 3.5 SIZE (nm}
05
(a) I
l
I
4.5
I
I
I
I
I
I
I
20
50 z bJ
g
z LtJ
k~
0 t.u
U_ W (D
~25
~D
n
oV 0.5
1.5
SIZE
25
\ F 0.8
3.5
/..5
(nm)
Fig. 9. Histogram of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 1.5 min at 650 °C.
2.0
I 32
I I I 4.; 5.6 6.B SIZE (nm)
B.O
9.2 10./.
(b) Fig. 10. Histograms of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 9.5 min at 6 5 0 °C.
198 40
I
I
distribution of particles can be markedly affected during the preparation of extraction replicas. A similar inhomogeneous particle size distribution is also observed after aging for about 60 min (Fig. 11). In addition, the frequency of the coarser particles (Fig. 1 l(b)) has also increased. Since it was not possible to estimate the particle volume fraction from extraction replicas [23], from the variation in the hardness of ferrite observed (Fig. 8) it can be inferred that the finer particles in the early stages of aging arise owing to nucleation with little growth and that an increasing particle volume fraction results in an increase in the hardness. In the later part of aging, however, a decrease in hardness occurs because of particle growth at constant particle volume fraction. The existence of finer particles at this stage can be explained by the particle dissolution and the growth of the more stable larger particles.
I
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5. Conclusions
Z IJA (21 W LI_
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(b) Fig. 11. Histograms of percentage particle frequency v s . alloy carbide precipitate particle size for steel D isothermally aged for 59.5 min at 650 °C.
about 10 min, the particle size distribution is observed to be inhomogeneous, with finer particles (Fig. 10(a)) mainly occurring within grains and coarser particles (Fig. 10(b)) at grain boundary regions. Particles near grain boundaries are coarser since early precipitation occurs in these regions, as the first transformation is initiated here. Hence these precipitates have an appreciably longer period for growth than those particles which nucleate at a later stage during aging. However, it must also be pointed out that the true
Alloy carbide precipitation during the isothermal aging of F e - C - V alloys at 650 °C starts simultaneously with austenlte transformation and occurs only in the ferrite phase. While the bulk of the carbide precipitation forms during the austenite-to-ferrite transformation, a significant proportion takes place after the finish of transformation. The kinetics of precipitation are strongly affected by alloying additions through their respective effect on the time-temperaturetransformation curves. Silicon speeds up the transformation, with the consequence that most vanadium carbide precipitation occurs within an aging time of 2 min. Silicon also increases the precipitation rate such that complete precipitation occurs after a further aging of about 3 min. On the contrary, manganese, by slowing down the austenite transformation, allows the bulk of the precipitation to occur after aging for 10 rain and complete precipitation after a further aging of about 10 rain. Moreover, other than affecting the austenite transformation behaviour, the addition of manganese has little influence on precipitation kinetics. The addition of manganese and silicon together results in an intermediate transformation rate such that bulk of the precipitation occurs after aging for about 7 min. The predominating effect of silicon alone in increasing the rate of
199
precipitation by about four times is little affected in steel D with combined addition. Acknowledgments The authors are grateful to the Science and Engineering Research Council for funding this work and to the Head of the Department of Metallurgy, University of Strathclyde, for providing research facilities. References 1 W. B. Morrison, J. Iron Steel Inst., London, 102 (1963) 317. 2 R. Phillips and J. A. Chapman, J. Iron Steel Inst., London, 204(1966) 615. 3 K. Bunghardt, K. Kind and W. Oelsen, Arch. Eisenhiittenwes., 26 (1956) 61. 4 R. K. Amin, M. Korchynsky and E B. Picketing, Met. Technol., 8 ( 1981 ) 250. 5 M. G. Akben, 1. Weiss and J. J. Jonas, Acta Metall., 29 (1981) 111. 6 T. N. Baker, Heat Treatment '73, Metals Society, London, 1975, p. 13. 7 W.B. Morrison, B. Mintz and R. C. Cochrane, Controlled Processing of High Strength Low Alloy Steels, Preprints BSC Conf., 1976.
8 W.G. Hall and T. N. Baker, Met. Sci., 15 ( 1981 ) 447. 9 N. A. Razik, G. W. Lorimer and N. Ridley, Acta Metall., 22 (1974) 1249. 10 G. R. Speich, Trans. Metall. Soc. AIME, 245 (1969) 2553. 11 F. E. Fujita and A. C. Damask, Acta Metall., 12 (1964) 331. 12 A. V. Cheremuslkina and M. I. Koraleva, Soy. Phys.-Solid State, 5 (1963) 330. 13 M. M. Yang and A. D. Krawitz, Metall. Trans. A, 15 (1984) 1545. 14 H. W. King and S. G. Glover, J. Iron Steel lnst., London, 188(1958)61. 15 S. Arajs, Phys. Status Solidi, 33 (1969) 731. 16 J. C. Swartz and L. J. Cuddy, J. Phys. Chem. Solids, 32 (1971) 385. 17 H. Waganblast and S. Arajs, J. Appl. Phys., 3 (1968) 5. 18 J. Takamura, I. Takahashi and M. Amano, Trans. Iron Steel Inst. Jpn., 9 (1969) 216. 19 P.J. Gavin, Scr. Metall., 3 (1969) 741. 20 W. Roberts and A. Sandberg, Report IM 1489, 1980 (Swedish Institute for Metals Research). 21 R. C. Sharma and G. R. Purdy, Metall. Trans., 4 (1973) 23O4. 22 T. Sakuma and R. W. K. Honeycombe, Met. Sci., 18 (1984) 449. 23 T. N. Baker and H. S. Ubhi, in G. W. Lorimer (ed.), Analytical Electron Microscopy, Proc. Electron Microscopy and Analysis Group, Institute of Metals, London, 1988, p. 135.