Bhadeshia H.K.D.H., Honeycombe R. Steels: Microstructure and Properties

88 CHAPTER 4 EFFECTS OF ALLOYING ELEMENTS ON FE–C ALLOYS

Fig. 4.15 Fe–0.75V–0.15C transformed 5 min at 725

◦

C. Interphase precipitation of VC in

ferrite (courtesy of Batte). Thin-foil electron micrograph.

Fig. 4.16 Fe–0.25V–0.05C transformed and held at 2½ h at 740

◦

C. VC precipitation on

dislocations (courtesy of Ballinger). Thin-foil electron micrograph.

4.4 STRUCTURAL CHANGES RESULTING FROM ALLOYING ADDITIONS 89

In general, the ﬁbrous morphology represents a closer approach to an equi-

librium structure so it is more predominant in steels which have transformed

slowly. In contrast, the interphase precipitation and dislocation nucleated struc-

tures occur more readily in rapidly transforming steels, where there is a high

driving force, e.g., in microalloyed steels (Chapter 10).

4.4.2 Alloy carbide fibres and laths

The clearest analogy with pearlite is found when the alloy carbide in lath morp-

hology forms nodules in association with ferrite. These pearlitic nodules are

often encountered at temperatures just below Ae

in steels which transform

relatively slowly. For example, these structures are obtained in chromium steels

with between 4 and 12 wt% chromium (Fig. 4.11), and the morphology is analo-

gous to that of cementitic pearlite. It is, however, different in detail because of

the different crystal structures of the possible carbides, e.g. Cr

is hexagonal

and Cr

is complex cubic. The structures observed are relatively coarse, but

ﬁner than pearlite formed under equivalent conditions, because of the need for

the partition of the alloying element, e.g. chromium between the carbide and

the ferrite. To achieve this, the interlamellar spacing must be substantially ﬁner

than in the equivalent iron–carbon case.

At lower temperatures the lath morphology is largely replaced by much ﬁner

ﬁbrous aggregates, e.g. in high Cr steels coarse laths of Cr

can be replaced

by ﬁne ﬁbres of the same carbide usually 500 Å in diameter. Their length, which

is determined by the size of the ferrite colony, can be up to 10 µm with little or

no branching. Similar morphologies occur,but are much less dominant, in steels

containing W, Ti, V and Nb.

Carbide ﬁbres are frequently associated with planar interfaces, as well as

with pearlitic-type interfaces. Nevertheless, these are boundaries which can

apparently propagate rapidly without the need for step migration. A computer

analysis of similar boundaries in austenitic steels has shown that they possess

comparatively high densities of coincident lattice sites.

4.4.3 Interphase precipitation

Interphase precipitation has been shown to nucleate periodically at the γ/α

interface during the transformation. The precipitate particles form in bands

which are closely parallel to the interface, and which follow the general direction

of the interface even when it changes direction sharply. A further characteristic

is the frequent development of only one of the possibleWidmanstätten variants,

e.g. VC platelets in a particular region are all only of one variant of the habit, i.e.

that in which the plates are most nearly parallel to the interface. The bands are

often associated with planar low-energy interfaces, and the interband spacing

is determined by the height of steps which move along the interface (Fig. 4.17).

90 CHAPTER 4 EFFECTS OF ALLOYING ELEMENTS ON FE–C ALLOYS

Fig. 4.17 Fe–12Cr–0.2C wt% transformed 30 min at 650

◦

C. Precipitation of M

at stepped

γ/α interface: (a) bright field; (b) precipitate spot dark field (courtesy of Campbell). Thin-foil

electron micrograph.

The nucleation of the carbide particles occurs normally on the low energy planar

interfaces, rather than on the rapidly moving high-energy steps.

The need for step movement on the γ/α interface is in contrast to the growth

of ﬁbrous carbides behind an interface on which no steps are observed. Indeed if,

in these circumstances, a step does move along the interface, the ﬁbrous growth

is stopped and replaced by interphase precipitation. The step height and, there-

fore, the band spacing of the precipitation, is dependent on the temperature of

transformation and on the composition.As the temperature of transformation is

lowered the band spacing is reduced, e.g. in a 1 wt% V 0.2 wt% carbon steel, the

spacing varies from 25 nm at 825

◦

C to 7.5 nm at 725

◦

C (Fig. 4.18), and at lower

temperatures has been observed to be less than 5 nm. The extremely ﬁne scale

of this phenomenon in vanadium steels, which also occurs in Ti and Nb steels, is

due to the rapid rate at which the γ/α transformation takes place. At the higher

transformation temperatures, the slower rate of reaction leads to coarser struc-

tures. Similarly, if the reaction is slowed down by addition of further alloying

elements, e.g. Ni and Mn, the precipitate dispersion coarsens. The scale of the

dispersion also varies from steel to steel, being coarsest in chromium, tungsten

and molybdenum steels where the reaction is relatively slow, and much ﬁner

in steels in which vanadium, niobium and titanium are the dominant alloying

elements and the transformation is rapid.

Ainsley, M. H., Cocks, G. J. and Miller, D. R., Metal Science 13, 20, 1979.

4.5 TRANSFORMATION DIAGRAMS FOR ALLOY STEELS 91

Fig. 4.18 Interphase precipitation of VC in vanadium steels. Precipitate sheet spacing as a

function of transformation temperature (courtesy of Ballinger).

4.4.4 Nucleation in supersaturated ferrite

It has been shown that ferrite can occur in different morphologies depending on

the transformation temperature. At the highest transformation temperatures,

equi-axed boundary allotriomorphs form at the austenite grain boundaries, and

carbon diffuses to the austenite. In alloy steels, e.g. lowV steels,there is evidence

that the alloying element can also partition. As a result no alloy carbide forms in

this ferrite, which is thus truly pro-eutectoid. At lower temperatures the ferrite

formed is still equi-axed, but the alloy carbide forms at the same time either

as interphase precipitate or as ﬁbres. This is probably the closest approach to

true eutectoid behaviour in an alloy steel containing a strong carbide-forming

element.

At still lower transformation temperatures the ferrite adopts a Widmanstät-

ten habit and forms as laths, as in pure iron–carbon alloys. However, this ferrite

can be supersaturated when ﬁrst formed. If held only for a short time at the

transformation temperature,precipitation of the alloy carbide occurs within the

ferrite on dislocations. Such behaviour would be expected in alloy steels with

acicular ferrite provided a strong carbide former such as V, Ti or Nb is present

although, in theory, similar structures should be possible in plain carbon steels.

4.5 TRANSFORMATION DIAGRAMS FOR ALLOY STEELS

The transformation of austenite below the euctectoid temperature can best be

presented in an isothermal transformation diagram (Chapter 3), in which the

92 CHAPTER 4 EFFECTS OF ALLOYING ELEMENTS ON FE–C ALLOYS

beginning and end of transformation is plotted as a function of temperature

and time. Such curves are known as time–temperature–transformation, or TTT,

curves and form one of the important sources of quantitative information for

the heat treatment of steels. In the simple case of an eutectoid plain carbon

steel, the curve is roughly ‘C’-shaped with the pearlite reaction occurring down

to the nose of the curve and a little beyond. At lower-temperatures bainite and

martensite form (see Chapters 5 and 6). The diagrams become more complex

for hypo- and hyper-eutectoid alloys as the ferrite or cementite reactions have

also to be represented by additional lines.

Alloying elements, on the whole, retard both the pro-eutectoid reactions and

the pearlite reaction, so that TTT curves for alloy steels are moved increasingly

to longer times as the alloy content is increased. Additionally, those elements

which expand the γ-ﬁeld depress the eutectoid temperature, with the result that

they also depress the position of the TTT curves relative to the temperature

axes. This behaviour is shown by steels containing manganese or nickel. For

example, in a 13Mn–0.8C wt% steel, pearlite can form at temperatures as low as

400 C. In contrast, elements which favour the ferrite phase raise the eutectoid

temperature and theTTT curves move correspondingly to higher temperatures.

The slowing down of the ferrite and pearlite reactions by alloying elements

enables these reactions to be more readily avoided during heat treatment, so

that the much stronger low-temperature phases such as bainite and martensite

can be obtained in the microstructure. The hard martensitic structure is only

obtained in plain carbon steels by water quenching from the austenitic condition

whereas, by the addition of alloying elements, a lower critical cooling rate is

needed to achieve this condition. Consequently, alloy steels allow hardening to

occur during oil quenching, or even on air cooling, if the TTT curve has been

sufﬁciently displaced to longer times.

FURTHER READING 93

Igarashi, M. and Swaragi, Y. Proceedings of the International Conference on Power

Engineering – 1997, Japan Society of Mechanical Engineers,Tokyo, Japan, p. 107, 1997.

Krauss, G., Steels: Heat Treatment and Processing Principles, ASM International, Ohio, USA,

1990.

Leslie, W. C.,The Physical Metallurgy of Steels, McGraw-Hill, Tokyo, Japan, 1981.

Marden, A. R. and Goldstein (eds) Phase Transformations in Ferrous Alloys, TMS-AIME,

Warrendale, 1984.

Pickering, F. B., Physical Metallurgy and the Design of Steels, Applied Science Publishers,

London, UK, 1978.

Pierson, H. O., Handbook of Refractory Carbides and Nitrides, Noyes Publications, USA,

1996.

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FORMATION OF MARTENSITE

5.1 INTRODUCTION

The quenching to room temperature of austenite in a steel can lead to the for-

mation of martensite, a very hard phase in which the carbon, formerly in solid

solution in the austenite, remains in solution in the new phase. Unlike ferrite or

pearlite, martensite forms by a deformation of the austenite lattice without any

diffusion of atoms. The deformation causes a change in the shape of the trans-

formed region, consisting of a large shear and a volume expansion. Martensite

is, therefore, often referred to as a diffusionless, shear transformation, which is

highly crystallographic in character because it is generated by a speciﬁc deform-

ation of the austenite. When the formation of martensite is constrained by its

surroundings, it forms as thin plates or laths in order to minimize the strain

energy due to the deformation.

5.2 GENERAL CHARACTERISTICS

The martensite reaction in steels normally occurs athermally, i.e. the fraction

transformed depends on the undercooling below a martensite-start temperature,

. The extent of transformation does not seem to depend on time,as expressed

in the Koistenen and Marburger equation

which describes the progress of

transformation below M

1 − V

′

= exp{β(M

− T

)} where β ≃−0.011 (5.1)

′

is the fraction of martensite and T

the temperature below M

, to which

the sample is cooled. This athermal character is a consequence of very rapid

nucleationand growth,sorapid that thetime takencan beneglected. Instead,the

Koistinen, D. P. and Marburger, R. E.,Acta Metallurgica 7, 59, 1959.

96 CHAPTER 5 FORMATION OF MARTENSITE

Table 5.1 The temperature M

at which martensite first

forms on cooling, and the approximate Vickers hardness of

the resulting martensite for a number of materials

Composition M

/K Hardness HV

ZrO

1200 1000

Fe–31Ni–0.23C wt% 83 300

Fe–34Ni–0.22C wt% <4 250

Fe–3Mn–2Si–0.4C wt% 493 600

Cu–15Al 253 200

Ar–40N

fraction transformed depends only on the number of nucleation sites triggered,

with the less potent sites contributing at higher undercoolings.

From Equation (5.1) it is evident that someaustenite remains untransformed

when T

is set to room temperature. This is referred to as retained austenite.

It is also clear that there is no martensite-ﬁnish temperature, M

but for conve-

nience, the latter is frequently deﬁned at the point where 95% of the martensitic

transformation is completed.

Martensite is not restricted to steels although its technological importance

in steels is unsurpassed. Table 5.1 lists a variety of materials which exhibit

martensitic transformation,together with M

temperatures and hardness values.

To obtain martensite, it is usually necessary for the steel to be cooled from

the austenite phase ﬁeld at a rate which is sufﬁciently fast to avoid all other

solid-state transformations such as ferrite and pearlite. This cooling rate can

be very high for plain carbon steels but quite slow for a heavily alloyed steel

containing large concentrations of austenite stabilizing solutes.

Martensite can form at very low temperatures, where diffusion, even of

interstitial atoms, is not conceivable over the time period of the experiment.

Table 5.1 gives typical values of the martensite-start temperature for a variety

of materials. Martensite plates can grow at speeds which are a signiﬁcant frac-

tion of the speed of sound in the steel, some 1100 m s

−1

. Such a high rate of

growth is inconsistent with diffusion during transformation. Transformations

which involve diffusion are much slower – the fastest recorded solidiﬁcation

rate is about 80 m s

−1

in pure nickel. The chemical composition of martensite

can be measured and shown to be identical to that of the parent austenite.

These observations demonstrate convincingly that martensitic transformations

are diffusionless.

5.2.1 The habit plane

The habit refers to the interface plane between austenite and martensite as mea-

sured on a macroscopic scale (Fig. 5.1). For unconstrained transformations this

5.2 GENERAL CHARACTERISTICS 97

Fig. 5.1 An illustration of the habit plane between austenite (γ) and martensite (α

′

Table 5.2 Habit plane indices for martensite. With the exception of

ǫ-martensite, the quoted indices are approximate because the habit planes

are in general irrational (the square root of 2 is not a rational number)

Composition/wt% Approximate habit plane indices

Low-alloy steels, Fe–28Ni {111}

Plate martensite in Fe–1.8C {295}

Fe–30Ni–0.3C {31510}

Fe–8Cr–1C {252}

ǫ-martensite in 18/8 stainless steel {111}

interface plane is ﬂat, but the need to minimize strains introduces some curva-

ture when the transformation is constrained by its surroundings. Nevertheless,

the macroscopic habit plane is identical for both cases, as illustrated in Fig. 5.1.

Steels of vastly different chemical composition can have martensite with

the same habit plane (Table 5.2), and indeed, other identical crystallographic

characteristics.

5.2.2 Orientation relationships

The formation of martensite involves the coordinated movement of atoms. It

follows that the austenite and martensite lattices will be intimately related. All

martensitic transformations therefore lead to a reproducible orientation rela-

tionship between the parent and product lattices. It is frequently the case that

a pair of corresponding close-packed

planes in the ferrite and austenite are

The body-centred cubic lattice does not have a close-packed plane but {011}

is the most

densely packed plane.

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