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

198 CHAPTER 9 THE TEMPERING OF MARTENSITE

Fig. 9.11 Tempering of an Fe–4Mo–0.2C steel. Thin-foil electron micrographs: (a) 4.5 h at

550

◦

C. Cementite transforming in situ to Mo

C and start of nucleation on dislocations (Raynor

et al., 1966). (b) 5 h at 600

◦

C. Mo

C precipitation on dislocations within a former martensite

lath (courtesy of Irani). (c) 30 min at 700

◦

C. Mo

C precipitation in ferrite laths inherited from

martensite M

C precipitation at lath boundaries (Irani).

identiﬁed by dark ﬁeld microscopy. On further tempering, the positions of the

original cementite particles are indicated by small necklaces of alloy carbides

which tend to be coarser than the matrix precipitation.

Figure 9.11a also illustrates the dislocation network characteristic of tem-

pered steels and inherited from the martensite, although there has been

considerable rearrangement and reduction in dislocation density. Dark ﬁeld

electron microscopy reveals that these dislocations are the sites for very ﬁne

precipitation of the appropriate alloy carbide. On further ageing the par-

ticles are more readily resolved, e.g. in Mo steels as a Widmanstätten array,

9.4 TEMPERING OF ALLOY STEELS 199

comprising Mo

C rodlets lying along 001

directions. Figure 9.11b illustrates

this stage in a single martensitic lath. Heavier precipitation is evident at the lath

boundaries.

The nucleation of carbides at the various types of boundary is to be expected

because these are energetically favourable sites which also provide paths for rel-

atively rapid diffusion of solute. Consequently theageing processis usuallymore

advanced in these regions and the precipitate is more massive (Fig. 9.11c). In

many alloy steels, the ﬁrst alloy carbide to form is not the ﬁnal equilibrium car-

bide and,in some steels, as many as three alloy carbides can form successively. In

these circumstances, the equilibrium alloy carbide frequently nucleates ﬁrst in

the grain boundaries, grows rapidly and eventually completely replaces theWid-

manstätten non-equilibrium carbide within the grains. This is illustrated in Fig.

9.11c for a 4 wt% Mo steel tempered 30 min at 700

◦

C, in which M

C equi-axed

particles are growing at the grain boundaries but Widmanstätten Mo

C is still

visible within the grains. It is interesting to note that the structure still possesses

the lath-shaped ferrite grains inherited from the martensite. Recrystallization

occurs after longer times at 700

◦

9.4.4 Tempering of steels containing vanadium

Vanadium is a strong carbide former and, in steel with as little as 0.1 wt% V, the

face-centred cubic vanadium carbide (VC) is formed. It is often not of stoichio-

metric composition, being frequently nearer V

, but with other elements in

solid solution within the carbide. Normally, this is the only vanadium carbide

formed in steels, so the structural changes during tempering of vanadium steels

are relatively simple.

Vanadium carbide forms as small platelets, initially less than 5 nm across and

not more than 1 nm thick. These form within the ferrite grains on dislocations

(Fig. 9.12a) in the range 550–650

◦

C,and produce a marked secondary-hardening

peak. There is a well-deﬁned orientation relationship (Baker/Nutting) with the

ferrite matrix: {100}

//{110}

. In the early stages of precipitation at 550

◦

the particles are coherent with the matrix, there being only a 3% misﬁt between

010

and 110

. However,at 700

◦

C,the platelets coarsen rapidly (Fig. 9.12b)

and begin to spheroidize. However, the original martensite laths can still be

recognized, and are only replaced by equi-axed ferrite grains after long periods

at 700

◦

Many steels containing vanadium, e.g. ½Cr½Mo¼V, 1Cr¼V, 3Cr1Mo¼V,

1Cr1Mo¾V,will exhibit extensive vanadiumcarbide precipitation on tempering,

because of the stability of this carbide, not only with respect to cementite but

also the several chromium carbides and molybdenum carbide (see Fig. 4.5).

Because of its ability to maintain a ﬁne carbide dispersion, even at temperatures

approaching 700

◦

C,vanadium is an important constituent of steels for elevated

temperature service.

200 CHAPTER 9 THE TEMPERING OF MARTENSITE

Fig. 9.12 Fe–1V–0.2C wt% quenched and tempered. Thin-foil electron micrographs: (a) 72 h

at 550

◦

C,VC nucleation on dislocations (courtesy of Raynor); (b) 50 h at 700

◦

C. Plates of VC

(courtesy of Irani).

9.4.5 Tempering of steels containing chromium

In chromium steels, two chromium carbides are very often encountered: Cr

(trigonal) and Cr

(complex cubic). The normal carbide sequence during

tempering is:

Matrix → (FeCr)

C → Cr

→ Cr

While this sequence occurs in higher-chromium steels, below about 7 wt%

Cr, Cr

is absent unless other metals such as molybdenum are present.

Chromium is a weaker carbide former than vanadium, which is illustrated by

the fact that Cr

does not normally occur until the chromium content of the

steel exceeds 1 wt% at a carbon level of about 0.2 wt%.

In steels up to 4 wt% Cr, the transformation from Fe

CtoCr

occurs

mainly by nucleation at the Fe

C/ferrite interfaces. Steels up to 9 wt% Cr do

not show secondary-hardening peaks in tempering curves (Fig. 9.13). However,

these curves do exhibit plateaus at the higher chromium contents, which are

associated with the precipitation of Cr

. Chromium diffuses more rapidly

in ferrite than most metallic alloying elements, with the result that Cr

detected during tempering at temperatures as low as 500

◦

C, and in compari-

son with vanadium carbide,chromium carbide coarsens rapidly. Thus, in a 2 wt%

Cr–0.2 wt% C steel, continuous softening will normally occur on tempering

between 500

◦

C and 700

◦

C, although addition of other alloying elements, e.g.

Mo, can reduce the rate of coarsening of Cr

In contrast, a 12 wt% Cr steel will exhibit secondary hardening in the

same temperature range (Fig. 9.13) due to precipitation of Cr

. Addition-

ally, Cr

nucleates at about the same time but at different sites, particularly

former austenite grain boundaries and at ferrite lath boundaries. This precipi-

tate grows at the expense of the Cr

which eventually disappears from the

9.4 TEMPERING OF ALLOY STEELS 201

Fig. 9.13 The effect of chromium on the tempering of a 0.35 wt% C steel (Bain and Paxton,

The Alloying Elements in Steel, 2nd edition, ASM, Ohio, USA, 1961).

microstructure, at which stage the steel has completely over-aged. This transi-

tion from Cr

to Cr

in high-chromium steels is by separate nucleation and

growth. Further alloying additions can promote one or other of these carbide

reactions, e.g. addition of tungsten encourages formation of Cr

by allowing

it to nucleate faster, while vanadium tends to stabilize Cr

. In doing so, it

decreases the rate of release into solution of chromium and carbon needed for

the growth of Cr

. Clearly, vanadium would be a preferred addition to tung-

sten, if a ﬁne stable chromium carbide dispersion is needed in the temperature

range 550–650

◦

9.4.6 Tempering of steels containing molybdenum and tungsten

When molybdenum or tungsten is the predominant alloying element in a steel,

a number of different carbide phases are possible, but for composition between

4 and 6 wt% of the element the carbide sequence is likely to be:

C → M

202 CHAPTER 9 THE TEMPERING OF MARTENSITE

Fig. 9.14 Fe–6W–0.23C wt% quenched and tempered (courtesy of Davenport).Thin-foil elec-

tron micrographs: (a) 100 h at 600

◦

C. W

C needles along 001

some irregular particles of

C. (b) 26 h at 700

◦

C. Massive M

The carbides responsible for the secondary hardening in both the case of tung-

sten and molybdenum are the isomorphous hexagonal carbides Mo

C andW

both of which, in contrast to vanadium carbide, have a well-deﬁned rodlet

morphology (Fig. 9.14a). When formed inthe matrix,M

C adopts aWidmanstät-

ten distribution lying along 100

directions. In molybdenum steels, peak

hardness occurs after about 25 h at 550

◦

C, when the rods are about 10–20 nm

long and 1–2 nm in diameter. The orientation relationship is:

(0001)

//(011)

[11

20]

//[100]

(rod growth direction).

C also nucleates at former austenite and ferrite lath boundaries. As in the

case of vanadium steels, M

C precipitate nucleates both on dislocations in the

ferrite, and at the Fe

C/ferrite interfaces, but the secondary hardening arises

primarily from the dislocation-nucleated dispersion of M

On prolonged tempering at 700

◦

C, the complex cubic M

C forms predom-

inantly at grain boundaries as massive particles which grow quickly, while the

C phase goes back into solution.The equilibrium microstructure is equi-axed

ferrite with coarse M

C in the form of faceted particles at grain boundaries, and

plates, illustrated in Fig. 9.14b for a 6 wt% tungsten steel tempered 26 h at 700

◦

For similar atomic concentrations, the secondary hardening response in the

case of tungsten steels is less than that of molybdenum steels. The M

C disper-

sion in the former case is coarser, probably because the slower diffusivity of

tungsten allows a coarsening of the dislocation network prior to being pinned

by the nucleation of M

C particles.

At lowerconcentration of tungsten and molybdenum (0.5–2 wt%), two other

alloy carbides are interposed in the precipitation sequences, i.e. the complex

9.4 TEMPERING OF ALLOY STEELS 203

cubic M

and the orthorhombic M

, probably Fe

MoC. These carbides

are found as intermediate precipitates between M

C and M

9.4.7 Complex alloy steels

The presence of more than one carbide-forming element can complicate the

precipitation processes during tempering. In general terms, the carbide phase

which is the most stable thermodynamically will predominate, but this assumes

that equilibrium is reached during tempering. This is clearly not so at tempera-

tures below 500–600

◦

C.The use of pseudo-binary diagrams for groups of steels,

e.g. Cr–V, Cr–Mo, can be a useful guide to carbide phases likely to form dur-

ing tempering (see Chapter 4, Section 4.2). The sequence of precipitation for

a particular composition can be approximated to by drawing a line from the

origin of the diagram, e.g. Fig. 4.6, to the composition of interest. The phase

ﬁelds passed through would normally be those encountered in tempering, but

the exact conditions cannot be forecast from such data.

Certain strong carbide formers, notably niobium, titanium and vanadium,

haveeffects on temperingout ofproportionto their concentration. Inconcentra-

tions of 0.1 wt% or less, provided the tempering temperature is high enough, i.e.

550–650

◦

C,they combine preferentially with part of the carbon and, in addition

to the major carbide phase, e.g. Cr

,Mo

C, they form a separate, very much

ﬁner dispersion, more resistant to over-ageing (Fig. 9.15). This secondary dis-

persion can greatly augment the secondary-hardening reaction, illustrating the

importance of these strong carbide-forming elements in achieving high strength

levels, not only at room temperature but also at elevated temperatures, where

creep resistance is often an essential requirement.

Fig. 9.15 Fe–4Mo–0.1Nb–0.2C wt% steel tempered 6 h at 700

◦

C. Coarse needles of Mo

in ferrite and fine particles of NbC on dislocations (courtesy of Irani). Thin-foil electron

micrograph.

204 CHAPTER 9 THE TEMPERING OF MARTENSITE

9.4.8 Mechanical properties of tempered alloy steels

A wide range of mechanical properties is obtainable by tempering alloy steels

between 200

◦

C and 700

◦

C. A typical example is shown in Fig. 9.16 for a steel

containing 1.5Ni–1Cr–0.25Mo–0.4C wt% (En24), the tensile strength of which

can be varied from 1800 down to 900 MN m

−2

by tempering at progressively

high temperatures. The ductility of the steel improves as the tensile strength

falls. However, there is a ductility minimum around 275–300

◦

C, which is often

observed in plain carbon and lower-alloy steels. This has been attributed to

the conversion of retained austenite to bainite, but it is more likely to be the

result of the formation of thin cementite ﬁlms, as a result of the transformation

of austenite at the interlath boundaries. At higher temperatures, these ﬁlms

spheroidize and the toughness improves.

To obtain really high strength levels in tempered steels (∼1500 MN m

−2

),it is

usual to temper at low temperatures, i.e. 200–300

◦

C,when the martensite is still

heavily dislocated and the main strengthening dispersion is cementite or ε-iron

Fig. 9.16 Mechanical properties of En 24 (1.5Ni–1Cr–0.25Mo–0.4C wt% steel) as a result of

tempering for 1 h (Thelning, Steel and its Heat Treatment, Bofors Handbook, Butterworth, UK,

1975).

9.4 TEMPERING OF ALLOY STEELS 205

Fig. 9.17 Comparison of mechanical properties of plain carbon and alloy steels tempered at

200

◦

C (Irving and Pickering, Journal of the Iron and Steel Institute 194, 137, 1960): (a) effect of C

on tensile strength; (b) relation between tensile strength and impact value. Note the beneficial

effect of Mo.

carbide. Alloy steels, when tempered in this range, not only provide very high

tensile strengths with some ductility but are also superior to plain carbon steels,

as shown in Fig. 9.17. It is clear from Fig. 9.17a that the carbon content has a large

inﬂuence on the strength. The alloying elements reﬁne the iron carbide disper-

sion and, as the carbon contentis raised,the dispersionbecomes moredense and,

therefore, more effective. The toughness decreases with increasing strength, as

206 CHAPTER 9 THE TEMPERING OF MARTENSITE

shown in Fig. 9.17b. However, alloying elements very substantially improve the

toughness, when compared with plain carbon steels of similar strength levels.

When molybdenum is present in the steel, the toughness is increased further as

the scatter bands indicate. This effect of alloying elements is again attributed to

the breakdown of carbide ﬁlms at grain and martensite lath boundaries. These

ﬁlms are particularly less noticeable in steels containing molybdenum.

Alloy steels which exhibit secondary hardening can provide high strength

levels on tempering between 500

◦

C and 700

◦

C, with better ductility than that

obtained at lower tempering temperatures. However, one of their main advan-

tages is that, once a high strength level is reached by means of an alloy carbide

dispersion formed between 550

◦

C and 650

◦

C, this structure will be relatively

stable at temperatures up to 500

◦

C. Therefore, the steels are suitable for use

under stress at elevated temperatures. A typical example is given in Fig. 9.18 of

Fig. 9.18 Effect of tempering for 1 h on the mechanical properties of a 12Cr–1Ni–0.2C wt%

stainless steel. Typical results for 50-mm diameter bars, oil-quenched (Thelning, Steel and its

Heat Treatment, Bofors Handbook, Butterworth, 1975).

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

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