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

288 CHAPTER 13 WELD MICROSTRUCTURES

Fig. 13.1 (a) Illustration of the epitaxial growth of columnar grains from the fusion boundary

of a stainless steel weld (courtesy of Honeycombe and Gooch). (b) Optical micrograph showing

the columnar prior-austenite grain structure typical in steel weld deposits.

and hence stiﬂethe growthof unsuitably orientedgrains.The widthof the colum-

nar grains therefore increases with distance away from the fusion boundary.

As already pointed out, the δ-ferrite undergoes a solid-state transformation

to austenite as the temperature decreases. The austenite nucleates at the δ–δ

grain boundaries and develops into a columnar austenite grain structure which

strongly resembles that of the original δ-grains (Fig. 13.1b).

The detailed shape and size of the austenite grains is of importance in the

evolution of the ﬁnal microstructure. The effect of the austenite grain size is two-

fold. Firstly, the number density of austenite grain boundary nucleation sites

changes inversely with the grain size. Coarse-grained weld deposits therefore

have a higher hardenability. The second, and more subtle effect, arises from the

columnar shape of the austenite grains, a shape which is like that of a hexagonal

prism. The grains are typically about 100 µm wide and about 5000 µm in length.

This is quite unlike an equi-axed grain structure, and because of the fewer grain

junctions involved, allows the hardenability of a weld to be larger than that of

a wrought alloy.

Solidiﬁcation does not occur under equilibrium conditions during weld-

ing. Solidiﬁcation-induced chemical segregation, and composition variations

due to uncontrolled changes in the welding conditions, make the solidiﬁca-

tion microstructure inhomogeneous. The amplitude of these variations becomes

larger as the alloy concentration increases.

13.2 THE FUSION ZONE 289

Table 13.1 A comparison of the chemical composition (wt%) of a submerged are weld

with that of the plate being welded, and the wire used as the consumable electrode.

The welding conditions used were 34V, 900Amp (D.C. positive), at a welding speed of

0.005 m s

−1

, with a calcium silicate flux.

CMnSiCuAlNO

Plate 0.21 1.0 0.2 0.05 0.04 0.01 0.004

Wire 0.14 1.5 0.2 0.31 0.01 0.01 0.001

Weld 0.16 1.1 0.3 0.16 0.01 0.01 0.053

Mineral ﬂuxes or inert gas shrouds are employed in order to protect the

hot metal against environmental attack during welding. Such protection is not

entirely effective,with the result that the oxide contentof weldstends tobe much

larger than that of wrought steel (Table 13.1). The oxide particles are entrapped

in the fusion zone during solidiﬁcation. As discussed later, these non-metallic

particles serve as heterogeneous nucleation sites and hence are of considerable

importance in the evolution of the microstructure. Table 13.1 reveals some other

interesting differences between the plate and weld compositions. The copper

concentration of the weld is large because in this case, the welding wire has a

copper coating to enable better electrical contact. The silicon concentration in

the weld is larger than both the wire and the plate, because the excess silicon is

acquired by decomposition of the protective ﬂux.These observations emphasize

the complexity of the welding process, in which the chemical composition of the

ﬁnal weld deposit depends on many variables, including the plate, wire and ﬂux

compositions.

13.2.2 The as-deposit microstructure

The microstructure obtained as the weld cools from the liquid phase to ambi-

ent temperature is called the as-deposited or primary microstructure. Its major

components include allotriomorphic ferrite, Widmanstätten ferrite, and acicu-

lar ferrite (Fig. 13.2). There may also be some martensite, retained austenite or

degenerate pearlite. These latter phases occur in very small fractions, and are

known by the collective term microphases. Bainite, consisting of sheaves of par-

allelplatelets,is not generallyfound in well-designedwelding alloys. Instead,aci-

cular ferrite is induced to nucleate heterogeneously on non-metallic inclusions.

In practice, the gap between the components to be joined has to be ﬁlled

by a sequence of several weld deposits. These multirun welds have a compli-

cated microstructure (Fig. 13.3). The deposition of each successive layer heat

treats the underlying microstructure. Some of the regions of original primary

microstructure are reheated to temperatures high enough to cause the reforma-

tion of austenite, which during the cooling part of the thermal cycle transforms

into a different microstructure. Other regions may simply be tempered by the

290 CHAPTER 13 WELD MICROSTRUCTURES

Fig. 13.2 (a) Schematic illustration of the essential constituents of the primary microstructure

in the columnar austenite grains of a steel weld deposit. (b) Scanning electron micrograph of

the primary microstructure of a steel weld (courtesy of Rees). The terms α, α

and α

refer

to allotriomorphic ferrite,Widmanstätten ferrite and acicular ferrite, respectively.

13.2 THE FUSION ZONE 291

Fig. 13.3 The macrostructure of a multirun weld, made by sequentially depositing a number

of beads in each of the 12 layers (courtesy of Reed).

deposition of subsequent runs. The microstructure of the reheated regions is

called the reheated or secondary microstructure.

13.2.3 Allotriomorphic ferrite

Allotriomorphic ferrite (α) is the ﬁrst phase to form on cooling the austenite

grains below the Ae

temperature. It nucleates at the columnar austenite grain

boundaries. Because these boundaries are easy diffusion paths, they become

decorated with thin, continuous layers of ferrite. The layers then thicken at a

rate which is controlled by the diffusion of carbon in the austenite ahead of the

transformation interface. Under isothermal conditions, the ferrite thickness S

changes parabolically with time t (Chapter 3):

S = α

1/2

, (13.1)

where α

is called the parabolic rate constant. This is illustrated in Fig. 13.4

for alloys with different carbon concentrations; note that the growth kinet-

ics become sensitive to the carbon concentration as the latter approaches the

solubility of carbon in the ferrite.

The magnitude of the parabolic rate constant depends on the equilibrium

compositions of the austenite and ferrite, and on the diffusivity of carbon in

292 CHAPTER 13 WELD MICROSTRUCTURES

Fig. 13.4 An illustration of the parabolic thickening of ferrite during isothermal transformation.

Each curve represents a Fe–1Mn–C wt% steel with the carbon concentration as indicated on

the diagram.

austenite (Chapter 3). Alloying elements such as manganese, which stabilize

austenite, are associated with a smaller value of α

. In welding, transformations

are not isothermal, but nevertheless, because nucleation is not rate limiting, the

fraction of allotriomorphicferrite obtainedcorrelates directly with the parabolic

rate constant (Fig. 13.5a).

The fact that the thickness of the ferrite varies with the square root of time,

means that the rate of growth decreases as the ferrite layer gets thicker. This

is because the distance over which carbon has to diffuse increases with time

(Fig. 13.5b). The growth rate for a given alloy goes through a maximum as a

function of temperature, because the driving force for transformation increases

with undercooling whereas the diffusivity decreases. Consequently, as the weld

cools to temperatures less than about 600

◦

C, the diffusional growth of ferrite

slows down so much that the layers of allotriomorphic ferrite reach a limiting

thickness. Widmanstätten ferrite formation does not involve the diffusion of

substitutional solutes, and therefore its growth is not sluggish at low tempera-

tures. The remaining austenite, therefore, begins to transform toWidmanstätten

ferrite (Fig. 13.6).

13.2.4 Widmanstätten ferrite and acicular ferrite

Although substitutional solutes and iron atoms do not diffuse during the growth

ofWidmanstätten ferrite, carbon does partition during transformation. Because

13.2 THE FUSION ZONE 293

Fig. 13.5 (a) The correlation between the calculated parabolic thickening rate constant

(a variable related to the growth rate) and the volume fraction of allotriomorphic ferrite

obtained in a series of manual metal arc weld deposits, fabricated using similar welding

parameters but with different chemical compositions. The rate constant is calculated for trans-

formation at 700

◦

C. (b) The diffusion distance increases as the ferrite layer thickens, slowing

down the rate of growth.

Fig. 13.6 Widmanstätten ferrite plates growing from allotriomorphic ferrite in a partially

transformed steel weld which was quenched from the transformation temperature.The matrix

is martensitic (courtesy of Barritte).

294 CHAPTER 13 WELD MICROSTRUCTURES

of its plate shape, much of the carbon can be accommodated at the sides of the

growing plate, so that the plate tip always encounters fresh austenite. This is

unlike the case for allotriomorphic ferrite, where the partitioned carbon builds

up ahead of the interface and progressively slows down the rate of growth.

Widmanstätten ferrite plates therefore lengthen at a constant rate.

The growth rates are found to be so large for typical weld compositions, that

the formation of Widmanstätten ferrite is usually completed within a fraction of

a second. Hence, for all practical purposes, the transformation can be regarded

as being isothermal (Fig. 13.7a).

Unfortunately, the fraction of Widmanstätten ferrite that forms in weld

deposits correlates badly with the plate lengthening rate, as illustrated in

Fig. 13.7b. This is because there is an interference between the plates of Wid-

manstätten ferrite that grow from the austenite grain boundaries, and acicular

ferrite plates which nucleate at non-metallic particles dispersed throughout the

weld (Fig. 13.8). The formation of Widmanstätten ferrite and acicular ferrite is

therefore competitive. Anything that increases the number density of inclusion

nucleation sites relative to austenitegrain nucleationsites,favours the formation

of acicular ferrite at the expense of Widmanstätten ferrite. Hence, the reﬁne-

ment of austenite grain size, or a reduction in the oxide content of the weld

below a limiting value, both lead to a decrease in the acicular ferrite content.

By the time the weld deposit cools to about 500

◦

C, most of the austenite

has been consumed. The small quantity of remaining austenite (about 5%) is

enriched in carbon and either transforms to martensite, or into pearlite, which

is degenerate because it does not have the opportunity to establish a lamel-

lar structure. Slower cooling rates favour the formation of pearlite relative

to martensite. Some austenite may also be retained to ambient temperature.

Because of their small volume fractions in the overall microstructure, these

phases are, in welding terminology, called ‘microphases’. The microphases are

relatively hard and behave in many respects like brittle inclusions. They are,

therefore, of importance in determining the toughness of weld deposits.

13.2.5 Sensitivity to carbon

It is striking that small variations in carbon concentration can have a major

inﬂuence on the microstructure of welds, especially since the average carbon

concentration of a weld is usually kept very small. It is apparent from the pre-

vious discussions of the growth rates of allotriomorphic and Widmanstätten

ferrite, that the sensitivity of growth kinetics to carbon becomes larger as the

concentration of carbon decreases.

These are important observations given that the general trend in the steel

industry is to reduce the carbon concentration, sometimes to levels approach-

ing the maximum solubility of carbon in ferrite. The rate at which ferrite

grows increases sharply as the carbon concentration of the steel approaches

13.2 THE FUSION ZONE 295

Fig. 13.7 (a) The isothermal growth rate of Widmanstätten ferrite in a series of

Fe–1Mn–C wt% alloys as a function of carbon concentration. Notice that the growth rates

are so large, that the plates could grow right across typical austenite grains within a fraction

of a second. (b) Poor correlation of the volume fraction of Widmanstätten ferrite against the

calculated growth rate.

its solubility in ferrite. This is because there is no need for the carbon to diffuse

ahead of the γ/α interface, since it can all be accommodated in the ferrite.

Hence, the effect of carbon is seen to be larger (Figs 13.4 and 13.6a) when its

concentration changes from 0.03 →0.05 wt%, when compared with the change

from 0.09 →0.11 wt%. Changes in mechanical properties are found to reﬂect

this behaviour, the strength of low-carbon steels being particularly sensitive to

296 CHAPTER 13 WELD MICROSTRUCTURES

Fig. 13.8 Diagrams illustrating the development of microstructure in two weld deposits with

different chemical compositions. The hexagons represent cross-sections of columnar austen-

ite grains whose boundaries first become decorated with uniform, polycrystalline layers of

allotriomorphic ferrite, followed by the formation of Widmanstätten ferrite. Depending on

the relative transformation rates of Widmanstätten ferrite and acicular ferrite, the former can

grow entirely across the austenite grains or become stifled by the intragranularly nucleated

plates of acicular ferrite.

the carbon concentration. This increased sensitivity of the γ/α transformation

to carbon at low concentrations, leads to a corresponding decreased sensitivity

to substitutional alloying elements. Carbon in effect controls the kinetics of

transformation.

In welding, the hardenability of the steel is often expressed as a carbon

equivalent (CE). The concentration of each solute is scaled by a coefﬁcient

which expresses its ability, relative to carbon, to retard the γ/α transformation.

Steels with a CE in excess of about 0.4 wt% cannot easily be welded because

of their increased tendency to form martensite. There are in fact two popular

expressions for the CE,one due to the International Institute forWelding (IIW),

and the other due toIto andBesseyo,covering the high and lowranges ofcarbon,

respectively:

IIW > 0.18 wt% C,

CE = C +

Mn + Si

Ni + Cu

Cr + Mo + V

wt%, (13.2)

Ito–Besseyo < 0.18wt% C,

13.2 THE FUSION ZONE 297

CE = C +

Mn + Cu + Cr

+ 5B wt%. (13.3)

The Ito–Besseyo CE formula has smaller coefﬁcients for the substitutional

solutes when compared with the IIW formula. It is believed to be more reli-

able for low-carbon steels. The IIW formula shows much smaller tolerance

to substitutional alloying elements than the Ito–Besseyo equation. As already

discussed, with low carbon concentrations the kinetics of transformation are

more sensitive to carbon than to substitutional solutes. Hence, it is logical that

there should be two different empirical expressions for the CE for the low- and

high-carbon weldable steels. Figure 13.9 illustrates that, as expected, both the

Fig. 13.9 Variations in microstructure and mechanical properties as a function of carbon

concentration in Fe–1Mn–C wt% steel weld deposit using manual metal arc welding (1 kJ/mm).

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

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