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

Подождите немного. Документ загружается.

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

Fig. 4.6 Carbide constitution in 0.2%C steels at 700

◦

C as a function of vanadium and

chromium content (Shaw and Quarrell, Journal of the Iron and Steel Institute 185, 10, 1957).

austenite occurs would be sensitive to the concentration of alloying elements in

steel. Both the growth of ferrite of pearlite are affected, so these reactions will

be considered separately. Most familiar alloying elements displace the time–

temperature–transformation (TTT) curve for a plain carbon steel to the right,

i.e. towards longer transformation times. However, a small group of elements

move the curve to shorter transformation times.

4.3.1 The effect of alloying elements on the ferrite reaction

Two basically different modes of growth of pre-eutectoid ferrite in austenite

have been observed in Fe–C–X alloys. The actual mode observed is depend-

ent on the composition of the alloy but the two modes may occur at different

temperatures in the same alloy (Chapter 3). The modes are:

(a) growth with partition of the alloying element X between α and γ under

local equilibrium conditions;

(b) growth with no partition of X between α and γ under local equilibrium

conditions.

In the ﬁrst mode, the ferrite grows at a slow rate determined by the diffusivity

of the alloying element X in the γ-phase. This behaviour is sensitive to alloy

4.3 EFFECT OF ALLOYING ELEMENTS ON THE γ/α TRANSFORMATION 79

Fig. 4.7 Effect of manganese and molybdenum on the kinetics of the ferrite reaction (Kinsman

and Aaronson, in Transformation and Hardenability in Steels, Climax Molybdenum Co., Michigan,

USA, 1967).

composition which is shown by the fact that an Fe – 1.3 at%, C – 3.2 at%, Mn

alloy exhibits Mn partition at 742

◦

C, whereas an Fe – 1.0 at%, C – 1.5 at%, Mn

alloy shows partition of manganese at 725

◦

C. Alloys in which X is Ni or Pt also

show partition at higher transformation temperatures.

The mode whereno partitionoccurs givesrise toanarrow zone ofenrichment

or depletion, depending on whether X is a γ-orα-stabilizer,which moves ahead

of the α/γ interface. Aaronson and Domian have shown the lack of substitutional

solute partitioning for alloys in which X =Si, Mo, Co, Al, Cr and Cu for all

temperatures investigated. Ni, Mn and Pt on the other hand showed a greater

tendency to partition to the γ-phase. In the no-partition regime the observed

growth rates are relatively high, being determined by the diffusivity of carbon

which diffuses several orders of magnitude faster than the metallic alloying

elements. However, it has been shown that in Fe–C–X alloys, ferrite still grows

much more slowly than in Fe–C alloys, even when no partition of X is observed.

Some of this retardation is because the substitutional solute affects the thermo-

dynamic stability of γ relative to α. This is illustrated in Fig. 4.7 from the work

of Kinsman and Aaronson for X =Mn and Mo. To explain the results for the

molybdenum containing alloy they proposed that the α/γ boundary collects

atoms during the transformation and, as a result, experiences an impurity drag.

A third approach to the ferrite reaction was introduced by Hultgren, who

proposed a state of para-equilibrium at the γ/α boundary. In this, the trans-

formation occurs at such a rate that the substitutional solutes are unable to

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

partition. Thus, the substitutional solute/iron atom ratio remains ﬁxed every-

where. Subject to this constraint, the carbon, which is a fast diffusing species in

iron, partitions to an extent which allows it to achieve local (para)equilibrium

at the interface. This is a metastable mode of transformation, which allows the

growth of ferrite to be controlled by the diffusion of carbon, without any par-

titioning of alloying element X. The latter primarily inﬂuences transformation

by affecting the thermodynamic stabilities of austenite and ferrite.

4.3.2 The effect of alloying elements on the pearlite reaction

The pearlite reaction is a typical nucleation and growth reaction and, under the

appropriate experimental conditions, rates of nucleation N and rates of growth

G can be determined (see Chapter 3). The work of Mehl and coworkers showed

that many alloying elements reduce both N andG. For example,in molybdenum

steels of eutectoid composition both N and G were decreased, and nickel steels

behaved in a similar manner. The growth rate G as a function of atomic concen-

tration of alloying elements in several groups of steels is shown in Fig. 4.8. The

change in slope for Mo steels was correlated with the substitution of cemen-

tite by a molybdenum-rich carbide. Certain elements, notably cobalt, increased

both N and G for the pearlite reaction. The rates of growth of pearlite nodules

at 660

◦

C in cobalt steels are compared with that of a Co-free steel in Fig. 4.9.

Fig. 4.8 Effect of alloying elements on the rate of growth of pearlite in the range 550–700

◦

(Mehl and Hagel, Progress in Metal Physics 6, 74, 1956).

4.3 EFFECT OF ALLOYING ELEMENTS ON THE γ/α TRANSFORMATION 81

Fig. 4.9 Effect of cobalt on pearlite growth rate (Mehl and Hagel, Progress in Metal Physics 6,

74, 1956).

Recent work on chromium steels has shown that the addition of 1 wt% Cr

to an eutectoid steel results in substantially lower growth rates of pearlite. It

follows that in general the C-curve for a pearlitic steel will be moved to longer

times as the concentration of alloying element is increased.

On examining the interface between pearlite and austenite during transfor-

mation,it appears that the basic nature of the pearlite reaction requires partition

of carbon between the cementite and the ferrite (Chapter 3). However, in the

presence of metallic alloying elements, it is not obvious, ab initio, whether par-

tition of these elements will take place, taking partition to mean partition at

the pearlite/austenite interface so that element X partitions between cemen-

tite and ferrite as they are formed. At a later stage of the reaction, and after

its completion, alloying elements can partition within the pearlite over a wide

temperature range.

It is now generally agreed that partition of X between cementite and ferrite

at the interface with austenite does occur in many systems, even at relatively low

transformation temperatures. While partition can be predicted on theoretical

grounds, it can now be investigated experimentally

using electron probe micro-

analysis, where a probe size of <0.1 µm allows the in situ analysis of pearlitic

ferrite and cementite in partly transformed alloys. Atom-probe techniques

Ridley N., Solid–Solid Phase Transformations, TMS-AIME, Pennsylvania, p. 807, 1981.

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

allow even higher resolution.

In this way the systems Fe–Mn–C and Fe–

Cr–C have been examined, and partition has been found at temperatures as

low as 550

◦

An approach to the pearlite reaction, similar to that described earlier for

the ferrite reaction, is to distinguish two modes of growth, a partition local

equilibrium and a non-partition local equilibrium situation,

which are both

temperature and composition dependent. Elements which favour the forma-

tion of austenite, and so depress the euctectoid temperature, and also have

low solubilities in cementite, e.g. Ni, will encourage the non-partitioning reac-

tion.Those elements which are strong ferrite formers and consequently raise the

eutectoid temperature, as well as being soluble in cementite,are likely to exhibit

the partitioning type of reaction at the higher transformation temperature, e.g.

Cr, Mo, Si. The growth of pearlite in the non-partitioning case is probably con-

trolled by volume diffusion of carbon in austenite, but this diffusivity is reduced

by the presence of other alloying elements, in part accounting for the observed

effect of elements such as Ni on the pearlite growth rate. Where partitioning

of X takes place, the diffusivity of the alloying element in austenite must be a

limiting factor.

Whatever the alloying element distribution is at the growing interface, sub-

sequent redistribution between the ferrite and the cementite takes place, i.e.

those elements with substantial solubility in cementite (carbide formers) will

diffuse into that phase and the non-carbide formers will not. In this way the

composition of cementite can vary over wide limits, e.g. manganese is very sol-

uble in Fe

C; up to 20% of the iron atoms can be replaced by chromium, while

vanadium will replace 10% and molybdenum only 4%. The change in compos-

ition of cementite, while not affecting the crystal structure, will inﬂuence, e.g.,

the pearlite interlamellar spacing, the detailed morphology and the tendency to

spheroidize.

Once the alloying element concentration reaches a critical level, the cemen-

tite will be replaced by another carbide phase. For example, in a chromium,

tungsten or molybdenum steel, the complex cubic M

carbide can form,

where M can include iron, chromium, molybdenum or tungsten (Figs 4.10 and

4.11). This change in the carbide phase does not necessarily alter the basic

pearlitic morphology and consequently alloy pearlites are obtained in which

an alloy carbide is associated with ferrite (Fig. 4.11). These pearlites occur

only in medium and highly alloyed steels, usually at the highest transformation

temperatures. At lower transformation temperatures in the same steel, cemen-

titic pearlite may still form because of the inadequate diffusion of the alloying

element.

Williams, P. R., Miller, M. K. and Smith, G. D. W., Solid–Solid Phase Transformations,TMS-

AIME, Pennsylvania, p. 813, 1981.

Coates, D. E., Metallurgical Transactions 4, 2313, 1973.

4.3 EFFECT OF ALLOYING ELEMENTS ON THE γ/α TRANSFORMATION 83

Fig. 4.10 Fe–12Cr–0.2C transformed 30 min at 775

◦

C. Pearlite-type reaction involving M

(courtesy of Campbell). Optical micrograph ×300.

Fig. 4.11 Fe–12Cr–0.2C transformed 15 min at 750

◦

C. Alloy pearlite. M

/ferrite (courtesy

of Campbell). Thin-foil electron micrograph.

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

4.4 STRUCTURAL CHANGES RESULTING FROM ALLOYING

ADDITIONS

The addition to iron–carbon alloys of elements such as nickel, silicon, man-

ganese, which do not form carbides in competition with cementite, does not

basically alter the microstructures formed after transformation. However,in the

case of strong carbide-forming elements such as molybdenum, chromium and

tungsten, cementite will be replaced by the appropriate alloy carbides, often at

relatively low alloying element concentrations. Still stronger carbide-forming

elements such as niobium, titanium and vanadium are capable of forming

alloy carbides preferentially at alloying concentrations less than 0.1 wt%. It

would, therefore,be expected that the microstructures of steels containing these

elements would be radically altered.

The tendency for forming carbides and nitrides can be expressed in terms of

bonding. Cottrell has been able to explain many of the observed trends in the

stability, crystal structure and stoichiometry of the carbides of transition metals

in terms of chemical bonds (see Further Reading). He points out that Ti, Zr

and Hf, which in the periodic table are elements near the beginning of the long

periods, form very stable MC carbides but the afﬁnity for carbon diminishes

further along the rows of the periodic table (Fig. 4.12). A part of the reason for

this is that more electrons have to be accommodated for elements further along

the rows,so antibonding statesare progressively ﬁlledthereby reducing thebond

order.

This does not completely explain the trend because the maximum bond

order occurs with Cr, Mo and W and we know that carbides of these elements

are less stable.

With MC carbides (where ‘M’ stands for metal atoms), the metal has to

sacriﬁce four electrons to form the bonds with carbon. Titanium has exactly

the right number so that its d-orbitals are left empty on forming TiC. This is

not the case with VC, since vanadium has an additional d-electron which forms

a V–V bond. The electrons in the two kinds of bonds, V–C and V–V mutually

repel, leading to a reduction in the stability of VC when compared withTiC.This

problem becomes larger along the row of the periodic table until MC carbide

formation becomes impossible or unlikely.

Although Cottrell has not considered the carbides in the lanthanide or

actinide series of elements, it is possible that the same principles should apply

there. Both NdCand UCexist. Remarkably,neodynium nitride hasalready been

incorporated into a ferritic creep-resistant steel by Igarashi and Sawaragi with

When two hydrogen atoms, each with a single electron, are brought together, they no longer

have separate atomic orbitals. Instead they have a pair of communal orbitals (bonding and

antibonding) each of which can hold two electrons. It follows that for H

both the electrons

are in the bonding orbitals giving a bond order of 2 and a strong molecule. For He

, on the

other hand, the four electrons ﬁll up both the bonding and the antibonding orbitals so the

bond order is zero, the molecule is not formed.

4.4 STRUCTURAL CHANGES RESULTING FROM ALLOYING ADDITIONS 85

Fig. 4.12 The periodic table showing the positions of strong carbide-forming elements.

rather good results. The concentration of neodynium used was only 0.04 wt%

but gave an increase in the creep rupture life by a factor of about two during tests

at 650

◦

C. They also tried hafnium but did not recommend it due to a tendency

to form coarse particles.

It has been shown how the difference in solubility of carbon in austenite

and ferrite leads to the familiar ferrite/cementite aggregates in plain carbon

steels. This means that, because the solubility in austenite is much greater than

in ferrite, it is possible to redistribute the cementite by holding the steel in

the austenite region to take it into solution, and then allowing transformation

to take place to ferrite and cementite. Examining the possible alloy carbides,

and nitrides, in the same way, shows that all the familiar ones are much less

soluble in austenite than is cementite. In Fig. 4.13 the solubility products in

austenite of vanadium, titanium and niobium carbides and nitrides are plotted

as a function of 1/T. Chromium and molybdenum carbides are not included, but

they are substantially more soluble in austenite than the other carbides. Detailed

consideration of such data, together with practical knowledge of alloy steel

behaviour, indicates that, for niobium and titanium, concentrations of greater

than about 0.25 wt% will form excess alloy carbides which cannot be dissolved

in austenite at the highest solution temperatures. With vanadium the limit is

higher at 1–2 wt%, and with molybdenum up to about 5 wt%. Chromium has a

much higher limit before complete solution of chromium carbide in austenite

becomes difﬁcult. This argument assumes that sufﬁcient carbon is present in the

steel to combine with the alloying element. If not, the excess metallic element

will go into solid solution both in the austenite and the ferrite.

4.4.1 Ferrite/alloy carbide aggregates

Steels containing strong carbide-forming elements transform from austen-

ite to ferrite in a similar way to, e.g., steels containing nickel or silicon.

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

Fig. 4.13 Solubility products of carbides and nitrides in austenite as a function of temperature

(Aronsson, in Steel Strengthening Mechanisms, Climax Molybdenum Co., Michigan, USA, 1969).

However, the carbide-forming elements restrict very substantially the γ-

loop (Fig. 4.4), so that the eutectoid composition is depressed to much

lower carbon levels and to higher transformation temperatures. One result

is that pearlite can completely disappear from the transformed microstruc-

tures, which now exhibit very different ferrite/carbide aggregates, usually on

a very much ﬁner scale than pearlite. Apart from the alloy carbide–pearlites,

particularly found in high chromium steels, there are three morphologies of

alloy carbides which are intimately associated with ferrite in the transfor-

mation temperature range in which plain carbon steels form ferrite/pearlite

structures.

4.4 STRUCTURAL CHANGES RESULTING FROM ALLOYING ADDITIONS 87

Fig. 4.14 Fe–4Mo–0.2C transformed 20 min at 650

◦

C. Fibrous Mo

C growth from

γ boundary (courtesy of Berry). Thin-foil electron micrograph.

Continuous growth of ﬁbres/laths The alloy carbides form as ﬁne ﬁbres or laths

which grow normal to the γ–α interface which then moves forward forming

ﬁbrous aggregates of carbide and ferrite (Fig. 4.14).

Repeated nucleation of carbides (interphase precipitation) In thisgrowth mode

the carbide particles, usually in the form of small platesor rods,nucleate atthe γ–

α interface which then moves to a new position where the nucleation cycle again

occurs. This process can be repeated many hundreds of time within a particular

austenite grain leading to a ferrite matrix with very ﬁne banded dispersions as,

e.g., in the 0.75 wt% vanadium steel shown in Fig. 4.15. Chromium steels give

coarser dispersions (Fig. 4.17).

Nucleation in supersaturated ferrite In microalloyed steels, where strong

carbide-forming elements are present in concentrations less than 0.1 wt%, it

is often possible to obtain the ferrite in a supersaturated condition with little

or no carbide precipitation taking place during the γ/α transformation. Instead,

while the steel is held at the transformation temperature, carbide precipitates

form within the newly formed ferrite grains, usually on dislocations (Fig. 4.16).

While it is possible by careful choice of alloy and experimental conditions

to obtain each of the above microstructures separately, in practice they are

often all present in transformed alloy steels, provided the steel contains a strong

carbide-forming element. Consequently the microstructures of transformable

alloy steels can be very complex, the full extent of these complexities only being

revealed when high-resolution electron microscopy is used to study them.