Transitions in carbide morphology in molybdenum and vanadium steels

METALLOGRAPHY 9, 277-291 0976)

277

Transitions in Carbide Morphology in Molybdenum and Vanadium Steels

A. BARBACKI Politechnika Poznanska (Technical University of Poznan), Physical Metallurgy Laboratory, Poland AND

R. W. K. HONEYCOMBE Department of Metallurgy and Materials Science, University of Cambridge, Pembroi;e Street, Cambridge UK

The decomposition of austenite to ferrite and alloy carbides has been investigated in a molybdenum and a vanadium steel by direct isothermal transformation at constant temperature and by changing the temperature during transformation. In this way transitions from interphase growth to fibrous growth of carbides have been studied using electron metallography. The experiments have also shown that the nature of the ~/a interface changes with transformation temperature, and it is concluded that this is the significant factor in determining the carbide morphology. The influence of crystallographic factors such as the degree of misfit between carbide and matrix are also discussed.

Introduction When alloyed austenite containing strong carbide-forming elements such as v a n a d i u m and m o l y b d e n u m is transformed to ferrite, alloy carbide dispersions can form [ 1 - 3 ] which considerably enhance the mechanical properties of the ferrite E3, 5]. While the resulting microstructures are frequently complex, two carbide morphologies tend to predominate. On the one hand, bands of fine carbide particles are formed, which have been shown to nucleate at the v / a interfaces [1, 2] as they m o v e through the alloy (interphase precipitation), the spacing of the bands and the particle size being dependent on the transformation t e m p e r a t u r e and other variables [5]. T h e second morphology is fibrous; in this the alloy carbides grow normal to the interpha~e boundary as fine fibers often only 500/~ diameter (~ American Elsevier Publishing Company, Inc., 1976

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A. Barbacki and R . W . K . Honeycombe

[4]. These two types of carbide morphology are competitive, often occurring together in the same specimen and even in the same original austenite grain. However, variations in heat treatment can markedly alter the proportion of the two morphologies observed. The purpose of the present investigation was to determine the factors which favor each morphology in steels of simple composition, and to examine the transition from one morphology to the other using heat treatments involving more than one transformation temperature. Two steels, one with molybdenum [-4] and one with vanadium [5], were chosen for the investigation because their normal transformation behaviour has received detailed study.

Experimental Details Two high purity steels were used: 1. 4.0%Mo 0.21%C 2. 0.98%V 1.58%Mn

0.025%Nb

0.18%C

all percentages being in wt%. Both alloys were produced by argon arc melting small ingots which were fabricated to 3-ram-diameter rod by hot and cold working. The specimens were sealed in silica tubes under reduced pressure of argon prior to heat treatment. Solution treatment was carried out at 1200°C for 80 min followed by water quenching. Prior to isothermal transformation, the specimens were reheated to 1200°C for 10 min then subjected to transformation in molten tin baths. °C f

850 800

!I

o;c

o+c

750

700

I

o.I

I

I.O

Fro. 1. TTT diagram of molybdenum steel with heat treatments adopted.

Morphology of Carbides

279

Results

1. T H E M O L Y B D E N U M S T E E L The T T T diagram determined for this steel is shown in Fig. 1. To investigate the formation of alloy carbides three heat treatments were used : (a) direct transformation at various temperatures in the range 725-850°C, (b) partial transformation at 725°C followed by further transformation at 850°C (Fig. 1, dash-dot lines), (c) partial transformation at 850°C followed by further transformation at 725°C (Fig. 1, dashed lines). The times of treatment were varied between 2 and 15 min to obtain between 10 and 20% transformation, after which the specimens were quenched to room temperature. 1.1 Direct transformation On transforming the steel directly at one temperature in the range 800-870°C the morphology changes in the following manner. At 870°C, the

Fro. 2. The 4% molybdenum steel: isothermal transformation at 850°C, showing interphase precipitation (A) and fibrous Mo~C. (B) The arrow indicates a step in the ~r/a interface. Optical micrograph. X 1600

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A. Barbacki and R.W.K. Honeycombe

carbide Mo~C is present in coarse bands of interphase precipitate which can be seen in the optical microscope. No fibrous carbide was observed at this temperature. However, on lowering the transformation temperature to 850°C, fibrous growth of Mo2C occurred in close association with coarse interphase precipitation (Fig. 2). It should be noticed that the latter forms in conjunction with coarsely faceted ferrite interfaces (arrow, Fig. 2), whereas the fibers grow at a smooth 7/a interface, confirming the observations of Campbell ~6] on chromium steels. 1.2 Decrease in temperature during the transformation (850 --* 725°C) Specimens partly transformed at 850°C were down quenched to 725°C where the transformation was allowed to proceed further before quenching into water (Fig. 1, dashed lines). Optical micrography revealed the replacement of some straight and stepped ferrite boundaries by curved boundaries which would be expected if interphase carbide precipitation were replaced by fibrous carbide growth at the lower temperature (Fig. 3). The transition is shown at higher magnification in Fig. 4, from an extraction replica, where the fine fibrous carbide morphology resulting at 725°C is

Fro. 3. 4% molybdenum steel: isothermal transformation at 850°C followed by isothermal transformation at 725°C. Change from interphase precipitation of Mo2C (A) to fibrous growth of Mo~C(B). Optical micrograph. X 600

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281

FI(;. 4. Same steel and treatment as Fig. 3. The arrow indicates a coarse step in the -~/a interface formed at 850°C. Extraction replica EM. X 6000 resolved. The arrow indicates a coarse step in the high-temperature ferrite boundary. Substantial regions of fibrous Mo2C formed at 850°C and on further transformation at 725°C, a fibrous structure was still formed b u t it comprised much finer, closer-spaced fibers as shown in Fig. 5, from an extraction replica of the transition zone. The change in spacing and diameter was approximately b y a factor of 10. The transition of interphase precipitate to fibers was also examined by thin-foil electron microscopy, which, when electron diffraction patterns were obtained from transition regions, revealed no changes in the orientation of the ferrite matrix, despite the marked change in carbide morphology. Figure 6(a) shows a typical transition region from coarse interphase precipitate on the left to fibrous Mo2C on the right. Precipitate spot dark field studies showed clearly that the carbides on each side of the transition possessed different orientations [Fig. 6(b)~, despite the fact that the ferrite matrix does not change orientation. Frequently more complex structures were observed, for example Fig. 7, in which the high-temperature transformation product (A) is in two zones, one coarsely fibrous and the other coarse interphase precipitation. Area B is the low-temperature (725°C) product in which the fine fibers take up a

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A. Barbacki and R.W.K. Honeycombe

FIG. 5. 4% molybdenum steel: isothermal transformation at 850°C then at 750°C. Transition from coarse fibrous Mo~C to fine fibrous Mo2C. Extraction replica EM. X 10,000 similar orientation to those in area A. However, a characteristic banding is imposed on these fine needles. Dark field photographs reveal that, while the needles have the same orientation, again the bands of interphase precipitation are of a different orientation. Electron diffraction photographs show t h a t all the ferrite in Fig. 7 is of the same orientation. 1.3 Increase in temperature during transformation (725 ° -~ 850°C) The structure usually formed at 725°C was fine fibrous Mo2C in ferrite (Fig. 8, area a). Up quenching to allow further transformation at 850°C led to large areas apparently devoid of carbide (b), which unlike low temperature ferrite (a), developed coarse growth steps. These regions had occasional patches of coarse fibers of Mo2C (c) and some coarse interphase precipitation of Mo2C (d). 2. T H E VANADIUM S T E E L 2.1 Direct transformation The TTT diagram of the vanadium steel determined by optical metallography is shown in Fig. 9. In the temperature range 700-760°C the

(a)

q

L_. . . . . . . . ..I

j

I

(b)

FIG. 6. Same steel and treatment as Fig. 5. (a) transition from coarse interphase precipitates (LHS) to fibrous Mo2C bright field E M , thin foil. (b) same area as (a). Precipitate spot dark field E M using diffraction spot from the fibrous carbide (RHS). X 50,000 283

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A. Barbacki and R.W.K. H oneycombe

Fro. 7. 4% molybdenum steel: isothermal transformation at 850°C then at 750°C. The high-temperature transformation product (A) comprises both coarse fibers and coarse interphase precipitate. The low-temperature region contains bands of fine Mo~C fibers (B). Thin-foil EM. X 40,000 austenite decomposes to a fine aggregate of vanadium carbide in ferrite, where the vanadium carbide is mainly as interphase precipitation, but occasional groups of fibers are also observed (Fig. 10). Frequently interphase precipitate regions also contained v e r y short fibers of vanadium carbide with the same orientation as t h a t of the interphase precipitate, which usually possessed a platelet morphology. 2.2 Increase in temperature during transformation (700 --~ 750°C) The type of treatment used is shown in Fig. 9 (dashed lines). A typical transition zone for such a treatment is shown in Fig. 11 in which area A is the lower-temperature zone. The morphology of the precipitate in zones A and B is similar except t h a t it is somewhat coarser in B, as would be

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285

F1G. 8. 4% molybdenum steel: isothermally transformed at 725°C then at 850°C. Area a contains fine fibrous Mo.2C formed at 725°C, while area b is ferrite with coarse growth steps. Fibrous Mo2C at c. Interphase precipitation at d. Optical micrograph. X 1600

°C

750

I' I

~

r------~

' J

¥~+¢ /

a+~

i

J

I

6OO J ©.1 FIG. 9.

I I,©

h

T T T diagram for the vanadium steel with the heat treatment adopted.

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A. Barbacki and R.W.K. Honeycombe

Fro. 10. 1% vanadium steel: isothermally transformed at 700°C. Interphase precipitation of VC also some small fibers of the same carbide. Thin-foil EM. × 30,000

FIG. 11. 1% vanadium steel" isothermally transformed at 700°C then at 750°C. A. 700°C zone, VC interphase precipitate. B. 750°C zone, VC interphase precipitate. Coarse VC at the changeover region. Extraction replica EM. X 40,000

Morphology of Carbides

287

expected from earlier work [-5]. However, the very coarse particles in the transition zone are unexpected; presumably they coarsen when the interface is stationary for some minutes when the specimen is first raised to 750°C (Fig. 9).

Discussion

The fact that interphase precipitation is normally associated with stepped ferrite interfaces, whereas fibrous carbides are more usually related to curved interfaces, has previously been explained in terms of the orientation relationship between the ferrite and the austenite grain in which it is growing [2]. The Kurdjumov-Saehs relationship is shown by Widmanst~itten ferrite growing in austenite and the planar lath boundaries of the ferrite are thus of low energy. In these circumstances, widening of the laths by the movement of incoherent steps would be expected, and has indeed been observed directly at temperature by means of thermionic emission microscopy [7]. It seemed thus logical to assume that interphase precipitation with its stepped ferrite interfaces was taking place when the ferrite was specifically orientation-related to the austenite in which it was growing. In contrast it was assumed that the fibrous carbide/ferrite aggregates behaved in a similar manner to pearlite, where the interface is normally curved and the ferrite is unrelated to the austenite in which it is growing [8]. The present experiments do not support this simplistic view, insofar as a change in transformation temperature of the molybdenum steel can alter the predominant carbide reaction from interphase to fibrous without any change in orientation of the ferritic matrix. This immediately indicates that the orientation relationship of the ferrite grain to the austenite in which it is growing is unimportant in determining the carbide morphology. Instead, there is evidence that the nature of the 7 / a interface is very important in determining the type of morphology adopted. The present work shows that, within a unique ferrite grain, as the transformation temperature is changed, the carbide morphology can alter, and with it the nature of the v/~ interface from stepped to curved, or vice versa. The experiments on the molybdenum steel were carried out at 850 and 725°C; referring to Fig. 1, to a first approximation the rate of this reaction is the same at these two temperatures, so we cannot ascribe the difference in the morphology of the carbide/ferrite aggregate to kinetic differences. The reactions both at 850 and 725°C are diffusion controlled, but at the lower temperatures the diffusivity of molybdenum in ferrite and in austenite will be reduced. In addition there is now increasing evidence that the

288

A. Barbacki and R.W.K. Honeycombe

interface boundary is an important diffusion path [-9], and that the diffusivity along a disordered boundary is greater than along a coherent or semicoherent low-energy boundary. We therefore propose that interphase precipitation gives way to fibrous carbide growth as the transformation temperature is lowered, because the 3'/a boundaries adopt incoherent high-energy configurations which allow more rapid diffusion along them, and hence a more rapid approach to equilibrium. At still lower temperatures (which have not been studied here) the ferrite grain grows with coherent interfaces (bainitic ferrite), but the reaction then is partly martensitic and not entirely diffusion controlled. The vanadium steel provides an interesting contrast with the molybdenum steel in so far as it exhibits predominantly interphase precipitation over the temperature range studied E5-], although in the present work this range was rather more limited than with the molybdenum steel. The interphase precipitation in the vanadium steel has usually been explained in terms of the rapid kinetics which such steels exhibit (Fig. 9). If the reaction is slowed down further by addition of more manganese, fibrous carbide occurs more extensively [-10-]; however, there is no clear evidence that in a given vanadium steel the amount of fibrous carbide increases with decreasing transformation temperature as it does for the molybdenum steel. Edmonds El01 has shown that even in a vanadium steel showing predominantly interphase precipitation there are very localized break-downs to fibrous carbide; however, in these cases the fibers are parallel to the original interphase carbide particles and arise simply by prolonged growth of these particles. This situation may arise if localized regions of the "~/~ boundary move not by the step mechanism along the interface, but by slow break away in a direction normal to the interface. This would allow time for fibrous growth which would stop when the step mechanism of growth was resumed. A further difference between the two steels is the crystallography of the two carbides, molybdenum carbide being hexagonal and vanadium carbide being fcc. Table 1 contains the lattice mismatches between the two carbides and both ferrite and austenite assuming the orientation relationships which have been observed. The misfit values of Table 1 are based on experimental evidence given by Berry et al. [-11, Harding [-111, and Dyson et al. [-122. Probable misfits for Mo~C growing in austenite were calculated on the basis of an orientation relationship common for the fcc --~ hcp transition, viz. (O0.1)h¢p I[ (111)~¢~ [ll.03rtep II ~llO-]f¢~

Morphology of Carbides

289 TABLE 1

Lattice Mismatch for Vanadium and Molybdenum Carbides in Ferrite and Austenite

VC Matrix

Orient. rel.

[100]~fl[ll0]~ [010]c]][110], (001)ell(001). y

cube/cube

Mo2C Misfit %

2.6 2.6 45.2 16.0 all directions

Orient. rel.

Misfit %

(00.1)ell(011), (01.0)~[l(011). (21.0)~11(100),

16.5 28.3 4.7

(00. l)~ll(lll)v (11.0)cl[(ll0)~ (11.0)c11(112)~

14.1 18.4 -11.2

The results in the table indicate more favorable crystallographic conditions for VC to grow in ferrite than in austenite, whereas with Mo2C the conditions are similar in both ferrite and austenite. Consequently it might be expected that ¥ C would grow more readily in ferrite, as it does in the case of interphase precipitation where it nucleates as single orientation plates as near parallel to the boundary as possible, presumably to maximize the effectiveness of the .y/a boundary diffusion path. The mechanism which has been proposed for this type of nucleation envisages the carbide platelets forming at the "7/a interface but growing entirely within the ferrite. The observed orientation relationship is a variant of {100}vc I[ {100}, (100)vc II (110), the Baker-Nutting relationship for vanadium carbide growing in tempered martensite (i.e., ferrite). In contrast, the growth direction of carbide fibers is approximately parallel to the direction of growth of the "y/a interface. The change in orientation of the Mo2C precipitate from needles, roughly parallel to the -y/a interface to fibers normal to the interface, is strikingly shown in Fig. 6(a). The interface must move sufficiently slowly to allow continuous fiber growth and in the case of the vanadium steel, this is achieved by addition of sufficient manganese or nickel, both of which displace the TTT curve to longer times. The greater preponderance of fibers in the molyb-

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A. Barbacki and R.W.K. Honeycombe

denum steel is probably in part due to the nature of the misfit between Mo2C and ferrite where there is a low degree of misfit in only one direction (see Table 1), in contrast to the plates of VC which have two directions of relatively low misfit. The other important factor is kinetic insofar as the -y -~ ~ ÷ carbide transformation is slower in the Mo steel, and consequently continuous fiber growth is encouraged. Conclusions 1. In both the vanadium and molybdenum steels interphase precipitation of VC and Mo~C is associated with facetted interfaces. 2. Fibrous carbide growth is associated with curved, high-energy ~/a interfaces. 3. Lowering of the temperature during the transformation can change the carbide morphology from interphase to fibrous, without a change in the ferrite orientation. 4. The changes in carbide morphology are accompanied by changes in the 7 / a interface from faeeted to high-energy boundaries. I t is concluded t h a t the nature of the ~ / a boundary determines to a large degree the precipitate morphology. 5. Crystallographic factors such as the degree of misfit between precipitate and matrix also influence the predominant morphology.

We would like to thank Prof. D.R. Miller, Dr. D.V. Edmonds, and Dr. P.A. Beaven for helpful discussions. One of us (A.B.) is also grateful for a British Council Visiting Scholarship. The work forms parl of a programme on alloy steels supported by the Science Research Council, U.K. References 1. F. G. Berry, A. T. Davenport, and R. W. K. Honeycombe, Institute of Metals Monograph No. 33, The mechanism of phase transformations in crystalline solids, 1969, p. 288. 2. A.T. Davenport and R. W. K. Honeycombe, Proc. Roy. Soc. A322, 191 (1971). 3. R. W. K. Honeycombe, in Effect of second phase particles on the mechanical properties of steel, ISI, 1971, p. 136. 4. F.G. Berry and R. W. K. Honeycombe, Met. Trans. 1, 3279 (1970). 5. A.D. Batte and R. W. K. Honeycombe, JISI 211~ 287 (1973). 6. K. Campbell and R. W. K. Honeycombe, Met. Sci. 8~ 197 (1974). 7. K. R. Kinsman, E. Eichen, and H. I. Aaronson, Met. Trans. 6A, 303 (1975). 8. R.J. Dippenaar and R. W. K. ttoneycombe, Proc. Roy. Soc. A333~ 455 (1973).

Morphology of Carbides

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9. D.A. Porter, D. B. Williams, and J. W. Edington, Proc. 8th International Congress on Electron Microscopy, Aust. Aead. of Science, Canberra, 1974, p. 656. 10, D.V. Edmonds, J!SI 210, 363 (1972). 11. H. J, Harding, Ph,D. Dissertation, University of Sheffield, 1966. 12. D, J. Dyson et el., Aeta Met. 14, 867 (1966).

Received July, 1975