CHAPTER-II

WELDING OF CR-MO-V STEELS FOR HIGH

STRENGTH & TOUGHNESS

2.1 Submerged Arc Welding (SAW)

2.2 The effect of Welding Parameters

2.3 Microstructure

2.4 Alloying Elements Addition

2.5 Effect of Carbides on Microstructures

References

CHAPTER -2

WELDING OF CR-MO-V STEELS FOR HIGH

STRENGTH VIS TOUGHNESS

Chapter I explain in general the hydrocracker reactors and the basic metal for its

construction. Chapter 2 deals with the various parameters of welding process and

metallurgical aspects of fabrication of hydrocracker reactors.

2.1 Submerged Arc Welding (SAW)

Submerged arc welding is an automated welding process in which the arc and

molten metal are both shielded from atmosphere by a blanket of fusible granular

material, reffered to as the ' FLUX '. Fig.2.1 schematically shows the SAW

process[4]. Since the arc is completely covered by the nux, it is not visible and

the weld is run without a nush or spark, and this is characteristic of open arc

welding. The advantages of the process include preventing the rapid escape of

heat due to the insulating blanket of nux deep penetration welds. The latter

characteristics means that the preparation for welding can be deep and narrow

with minimum use of filler metal [7].

The weld bead appearance is generally uniform due to the automation involved in

the process and hence mechanical properties in the weld tend to be at least equal

to those of the base metal. Also, very high deposition rates are possible in

SAW[9]. There are different types of consumables available for SAW. These

include single pass and multipass consumables. Multipass welds should be

deposited when welding carried out to ensure the specified heat input not

exceeded. A preheat of ISO°C is recommended for plate exceeding 16 mm in

thickness. Basically recommended that the root run for SAW can be deposited by

MMA Wand then ground back to base metal.

31

Residual nux layer

EJectrode wire

Welding head I AC or DC current Jaws ~ ~ ~ supply lead

~ Flux feed from hopper

Earth lead ..... ~ ¢,:; connection ............

"", ---C"I'/o.., ~ Flux layer 0, ... ",1.:

0'4) '9

Figure 2.1 Submerged Arc Welding[7J

The wire is coiled on a reel or in a drum. The nux is deposited on the joint ahead

of the arc from a flux hopper through a flux delivery tubc. After the weld metals

have solidified. the unfused flux is removal manually or by a vacuum pick up

system, \0 be sGcencd and reused. The weld backing plate severs to prevent burn

through; starting and run off tabs eliminate defects likely to occur during the stmt

and !Inish of a weld runtS].

2.2 The Effect of Welding Parameter

In SAW the weld deposit quality is determine by the type of nux. grade of wire

and the following parameter[ 14) :

a) Welding current

b) Arc Voltage

c) Speed of arc travel

d) Size of electrode

e) Electrode stick out

t) Heat input rate

32

To get optimum results, one must know their effect and how to select and control

them proper! y.

2.2.1 Welding Current[15]

It controls the melting rate of the electrode and thereby the weld deposition rate. It

also controls the depth of penetration and thereby the extent of dilution of the

weld metal by the base metal. Too high a current causes excessive weld

reinforcement which is wasteful, and burn through in the case of thinner plates or

in badly fitted joint, which are not provided with proper backing[5,7]. Excessive

current also produce a high narrow bead and under cut. Excessively low current

gives an unstable arc, inadequate penetration and overlapping. SAW equipment is

usually provided with an ammeter to monitor and control the welding current.

2.2.2 Arc voltage

Arc voltage, also called welding voltage, means the electrical potential difference

between the electrode wire tip and the surface of the molten weld puddle. It is

indicated by the voltmeter provided on the equipment[ 19]. It hardly affects the

electrode-melting rate, but it determines the profile and surface appearance of the

weld bead. As arc voltage increases, the weld bead become wider and flatter, and

the penetration decreases[ 17]. The effects of changing voltage are explained as

following:

I) Increasing voltage

a) Produces a flatter and wider bead

b) Increases flux consumption

c) Increases resistance to porosity caused by rust or scale

d) Helps to bridge gaps when fit up is poor

33

e) Increases pick out of alloy from the flux : this can be used to

advantages to raise the alloy content of the weld when welding is

performed with alloy or hard surfacing fluxes. If excessive, it can

reduce ductility and increases crack sensitivity, particularly in the

multiple pass welds[ 14].

2) Excessively high voltage:

a) Produce a hat shaped bead that is subject to cracking

b) Produces poor slag removal in groove welds

c) In multiple pass welds, increases the normal alloy pickup from the

flux

d) Produces a concave fillet weld that may be subject to cracking

e) Increases undercut on fillet welds

3) Lowering the voltage produces a stiffer arc needed for getting

penetration In a deep groove and to resist arc blow on high speed

work[ 12].

4) An excessively low voltage produces a high, narrow bead with poor

slag removal.

2.2.3 Speed of Arc Travel

For a given combination of welding current and voltage, increases in the welding

speed or the speed of arc travel results in lesser penetration, lesser weld

reinforcement and lower heat input per unit length of weld. Excessively high

travel speeds decrease fusion between the weld deposit and the parent metal, and

increases tendencies for undercut, arc blow, porosity and in'egular bead shaperS].

As the travel speed is decreased penetration, because under this condition, the

34

weld puddle is directly under the electrode tip and the force of the arc is

cushioned by the weld puddle[ 19]. Excessively slow speed also produces a

convex bead shape that is amenable to cracking, cause the arc to break out of the

flux layer which can hurt the operators eye, and give a large uncontrolled weld

puddle, which result is an uneven weld bead with slag inclusions.

2.2.4 Electrode Stick Out

It is also termed electrode extension. It refers to the length of the electrode,

between the end of contact tube the arc, which is subject to resistance heating

(also called eR heating) at the high current densities used in the process. The

longer the stick out, the greater the amount of heating and the higher the

deposition rater 16].

Increased electrode stick out, reduce to some extent the energy supplied to the are,

resulting in lower arc voltage and a different bead shape. Hence when the

electrode sticks out is increased to obtain higher deposition rate, the voltage

setting on the equipment must be increased to maintained correct arc length[18].

However with longer stick out, the increased in deposition rate is accompanied by

a decreases in penetration. Hence longest sticks out must be avoided when deep

penetration is desired. On the contrary, it can be used with advantages when melt

through is a problem while welding thin section[ 19]. It can also be used to

advantage in build up and hard facing application. It must be borne in mind that

with long electrode stick outs, it becomes more difficult to maintain the electrode

tip correctly in line with the joint.

35

Maximum electrode sticks outs recommended are;

75 mm for 2.0,2.4 and 3.2 mm wire diameter

125 mm for 4.0,4.8 and 5.6 wire diameter

2.2.5 Heat Input Rate (HIR)

Also termed arc energy, it is calculated by using the formula [7,5];

HIR= V X A X 60

S X 1000

Where, HIR = Heat input rate in kilojoules per mm

V = Arc voltage

A = Welding current

S = Arc travel speed in mm/min.

For a given joint thickness, the higher the heat input rate, the lower is the cooling

rate of the weld metal at heat affected zone (HAZ) of the parent metal, and vice

versa[3]. Heat input rate has an important bearing on the weld metal

microstructural and the final microstructure of the HAZ, and thereby on their

toughness. The significance of the parameter will be better understood while we

discuss the relation between SAW procedures and weld deposit toughness.

36

2.2.6 Submerged Arc Welding Problems[7]

The SAW process encounters following problems:

I. Solidification Cracking

2. Hydrogen Cracking

3. Incomplete fusion

4. Irregular wire feed

5. Porosity

2.2.7 Advantages and Application! 12]

Because of the following merits, the process of submerged welding is extensively

used in heavy steel plate fabrication.

I. Absence of smoke and arc flash

2. High weld metal quality

3. Smooth and uniform weld finish with no spatter

4. Extremely high deposition rate and welding speed

5. High electrode deposition efficiency

6. lIigh arc time can be achieved through automation

7. Welder's manipulative skill not needed; minimum operator fatigue.

S. Application cover pressure vessels, line pipe, storage tanks, heavy

structure, ships, railway wagons and coaches, and surfacing and build-up

work.

37

2.3 Microstructure

2.3.1 General

High strength steel weld metals with high yield strengths are available.

Microstructures suitable to resist stresses of such magnitude, are usually

composed of, or a combination of, martensite, bainite, and retained austenite .

Prior to these final constituents, the weld metal will have undergone complex

changes and transformations as it cools from its liquid form when deposited[ 12].

This may be clearly secn in the basic phase diagram of steel showing the iron -

carbon system as presented in Figure no 2.2[15].

However, one must bear in mind, that there are a few characteristic differences

between the microstructure of weld metals and those of wrought steels. Once

wrought steel is formed, various types of heat treatment may be performed in

order to adjust both the microstructure and the mechanical properties. However,

with weld metals this is not normally the case, and in many circumstances it is not

possible. Weld metals are cast materials and their cooling conditions are very

different in comparison to that of a steel that has undergone forging or rolling.

Other characteristic differences between two are the frequent presence of non­

metallic inclusions within weld metals and also segregation[24j. The occurrence

of interdendritic segregation must always be kept in mind when investigating the

weld metal microstructure. Basically there is a difference in composition between

the core and the other portions of the dendrite that formed during solidification.

38

P" • i ! ~ ...

0 20 1800

1700

1600

A l_

0

H 6 1400 -

N

1300 l_-T

1200 O~ T

1100 E n47- e

1000

,.FeJC

723" K-

&OO~~O~5--~I~~-f'5~~20~-d25~-J~O~-J~5~-'~O~~'~5~~S~ W.....," Con>on

Fig. 2.2 The Iron Carbon Phase Diagram[17]

As alloying is increased, it plays a more important role in determining the

microstructural transformations that take place and it may lead to different phases

in the dendrite-core and in interdendritic regions. Microsegregation allows the

formation of microstructures that are hard to produce in any other metallurgical

process. This can be beneficial when it is properly controlled but can create a lot

of difficulties for example crackingl161. When weld metals are deposited in multi

runs or if they are heat treated. part of the previous bead is reheated to a high

temperature and the bead is no longer in the "as deposited·· condition.

39

Microsegregation can be partially removed, however, a long annealing time IS

needed to remove it completely.

A brief overview of some phase transformation mechanisms that occur in high

strength steel weld metals will now be given; comprehensive treatments of the

physical metallurgy may be found in literature.

Introduction of Phase Diagram

As the properties of a material depends to a large extent on the number, amount,

type and the form (shape and size) of the phases present, and can be altered by

changing these quantities. Thus, the study of phase relationship plays an

important role in the understanding of the properties of the materials. Actually,

the state of a material depends on the variables/conditions like composition,

temperature and pressure[23]. The best way to understand these effects is through

phase diagrams, or equilibrium diagram. The equilibrium diagram is a map which

gives relationship between phases in equilibrium in a system as a function of

temperature, pressure and composition. It gives a "blue print" of the alloy system

from which we can anticipate at what compositions the alloys are likely to have

useful properties, what treatments must be given to alloys to develop their

properties to the best effects, and what treatments are likely to be harmful and

must be avoided.

Equilibrium diagrams also called constitutional diagrams illustrate the stable

states of a metal, or alloys[22J. i.e., those phases which have lowest free energy

under given conditions, which means phases obtained under very slow cooling, or

heating rate. When graphical representation deals with phases which are in

equilibrium with the surroundings, it is called an equilibrium diagrams otherwise,

it is called a phase diagrams.

40

Tn the iron carbon system, stable phases are solid solution of carbon in iron and

graphite. However under normal conditions, the solid solution and the compound,

Fe,C (cementite) are fOfll1ed[18). The graphical representation of Fe-Fe,C

system is a phase diagram (or a metastable diagram), while that of iron-graphite is

an equilibrium diagram.

2.3.2 Phase transformation

There are several phase transformations during the cooling of steel from the liquid

state to room temperature. One of the most important in welded plain carbon

steel, in terms of mechanical properties, is the solid-state transformation y --+ a +

Fe,C[ 14). This transformation occurs as a eutectoid decomposition at 723°C

(designated AI) under quasi-equilibrium conditions. Pearlite is the name given to

the eutectoid mixture of lamellar <x(ferrite) and Fe,C(cementite). As a result of the

hypo elltectoid C content of steel, proelltectoid ferrite is likely to form on slow

cooling from austenite, initiating at the Ar, temperature.

The presence of the strong carbide former Cr and Mo modify the composition and

structure of the carbide phase. The alloying elements also affect the temperature

quoted at 780'C in steel, and the AC3 temperature at 890°C and 880°C with

unspecified heating rater 16).

Under continuous cooling conditions such as those experienced during a weld

thermal cycle, the cooling rate determines the deviation from quasi-equilibrium

and therefore the morphology and extent of proeutectoid ferrite and carbide

precipitation[IS). The product that form are metastable since they are produced

under non-equilibrium conditions.

41

With faster cooling, the distribution of the carbide phase will change, as will the

morphology of the ferrite. Austenite transforms to a bainitic type product at

moderately fast cooling rates. Bainite is a eutectoidal product of medium to high

carbon steel, which contains lathe of ferrite with the carbide oriented along these

lathe boundaries(upper bainite) or distributed within the ferrite lathe(lower

bainite)[16]. Typically, multiple lathe form together with a common orientation,

forming a 'sheaf or 'packet'. The type of bainite formed depends on the

transformation temperature and therefore cooling rate.

Under isothermal transformation conditions, upper bainite typically forms in steel

between 400 and 550°C and lower bainite between 250 and 400°C. The kinetics

of the reactions are dependent on the composition of the steel.

The bainitic structures that form in low carbon steel are often described as

'bainitic ferrite' and 'granular bainite'. These microstmctures display bainitic

ferrite sheaves, without the carbides found in conventional bainite[24]. On

continuous cooling, the austenite transforms to bainite sheaves, rejecting carbon

and enriching and stabilizing the remaining austenite in the process.

This enriched austenite does not transform to carbide as in the conventional upper

bainite reaction[24]. Instead the austenite is retained to low temperatures where it

may either remain in its metastable state or transform fully or partially to one of

several 'residual phases'.

In steel, the bainite product is likely to lie between the two extremes of the

conventional upper/lower bainite of high carbon steel and the bainitic

ferrite/granular bainite of low carbon steel. Due to a carbon level of 0.15% and

the presence of strong carbide formers, limited carbide precipitation is to be

expected depending on cooling rate[12].

42

The alloying elements present in steels and other low alloy grades retard the

ferrite and pearlite reactions and depress the start temperature of the bainite(538-

566°C) and re-softening has been observed 100 hrs. These particles initially

precipitate as acicular, elongated needles.

Subsequent time at temperature sees the hardness decrease at a steady rate. This

softening is attributed to the coarsening and spheroidising of the Moze needles

and an increase in alloy content within the carbide particles.

2.3.2 CCT Diagram

Result compiling Table 2.1 and Table 2.2 shows the effect of tempering and

cooling rate of the test piece[181.

Table 2.1 Formation of different micro structure at different cooling rate.

Sr Cooing rate Ferrite Bainite Pertite Martensite

1 5°C/min ----- 100% ----- -----

2 12°C/min 15% 77-80 % ----- 5-8%

3 150°C/min 10% 20% 20% 50%

4 650°C/min ----- ----- ----- 100%

Table 2.2 Effect of cooling rate on hardness at different Bar Dia

Bar Dia. Sr Cooing rate Hardness Air Oit Water

1 5°C/min 40 HRC 500 500 500

2 12°C/min 42 HRC 200 500 500

3 150°C/min 45 HRC 20 150 180

4 650°C/min 52.5 HRC 5 65 85

Table 2.1 clearly indicates the formation of Bainite at slow cooling rate. The

highest cooling rate forms 100% Martensite.

Table 2.2 correlates the same data with different quenching media again the effect

of tempering is predominantly exhibited showing the propogation of tempering

43

effect at varIOUS diameter. It can be concluded that formation of Martensite IS

'C

directly proportional to cooling rate hence tempering effect.

~l, '"-'0 i \< ,

I ':'i)o)~

,orr

J

2~ Cr Mo V

,

+-T

-n:r, '. ~

\t '~11 '-it n~Ll> .\ I '~II, ( f'RI-YltltJ."i IKt-·\I,\1u,n ROULI) ",,1 :, .... '1) --'I

'" ,\I

'" "11>,

• I -----+ --

,. ; .' '- tlr ': ~ _ .~+

1-­DA~ I.

_. - , 1 I ".lI~

DlA'1FTFRr .','", •

"~I ,,~,

I ---__ 1 ;. """ ~ I :/.

, .. ,

I ,.".L _______________ -L ____ -'-_________ -'

Fig 2.3 CCT Diagram

44

2.3.4 Stress Relief Cracking

This form of cracking occurs as a weldment is reheated either in service or as part

of a post weld heat treatment. For this reason. it is also referred to as reheat

cracking. The cracking primarily occurs in the IIAZ and is intergranular with

respect to the prior austenite grain boundaries[ 12].

There are two forms of stress relief cracking, operating at different temperatures.

The higher temperature process involves fracture due to micro void coalescence.

This mechanism is initiated by the presence of particles at grain boundaries.

These can be compounds of contaminant species or of alloying elements such as

cm·bides.

These particles can deeohere from the matrix, effectively causing a discontinuity

or void in the matrix. The composition of these particles influences the cohesive

strength at the matrix/particle interface. This cracking mechanism is sometimes

referred to as embrittlement and can occur after long period of service or after

secondary thermal treatments[ 15].

The second. lower temperature form of stress relicf cracking is characterized by

low ductility intergranular cracking. The mechanism for this form of crack

development is related to the relative strengths of the grains and grain boundaries.

The precipitation of a particle dispersion within the bulk grain stmcture gives an

alloy increased strength[ 16]. Grain boundaries act as favoured nucleation and

growth regions and can contain coarsened carbides surrounded by precipitate free

zones (PFZ). As such, the grain boundary regions are less effectively strengthened

by precipitates and therefore weaker.

45

This strength difference within a grain causes localized strain to build up at grain

boundaries and the nucleation of cracks when exposed to certain loading and

thermal conditions. The initiation of this cracking is dependent on the existence of

a triaxial stress state. Such a stress state is observed in the residual stress fields of

welded joints as discussed above[27].

Stress relief cracking by both mechanisms has been reported for the 2.2SCr-IMo­

O.2SV alloy system[13]. The effect of various levels and types of impurity species

have also been investigated in these studies, with sulphur, tin and phosphorus all

aiding in these low ductility fractures. There is also a synergistic effect when

several speCIes are present, such as in order, low purity Cr-Mo commercial

alloys.

2.3.5 Solidification of Weld

When welding in particular material with similar or the same composition as the

molten weld metal, epitaxial growth rather than nucleation occurs. Epitaxial

growth is where the molten liquid solidifies and develops new grains directly from

the solid in the underlying material. Solidification generally occurs along the

maximum thermal gradient towards the weld centre resulting in a columnar

solidification structure[ 19]. The width of the columns are usually related to the

size of the grains in the underlying bead. It is nonnally expected that the grains

increase in size and are expected to be wider than the underlying grains. Grains

separated by high angle boundaries are called the primary grains and it is usually

found that there is a substructure divided by low angle grain boundaries within the . .

pnmary grams.

The substructure develops in different ways depending on the solute content, the

thermal gradient and the solidification velocity. Four different modes are usually

46

found; planar front, cellular, cellular dendritic and dendritic. Growth occurs from

the motion of an unstable solidification front and dendrites develop sometimes

with a slight deviation from the maximum thermal gradient direction following

the easy growth direction of <I 00>.

The Microstructural transformations that take place are largely dependant on both

the cooling ratc and the composition. In steel, it was reported that austenite

transformed to ferrite on cooling between 910 and 723°C in iron - carbon alloys.

At fast cooling, the transformation temperature may be depressed below the

rcgion of 690°C with austcnite transforming to upper bainite. At even higher

cooling rates transformation can take place in the rcgion of 500 °C or lowcr, i.e.

the martensite start temperature (Ms) (that depends on alloying). The three major

phases arc discussed in detail.

2.3.6 Ferrite

There are three main ferritic constituents that form in weld metals; allotrimorphic

ferrite, Widmanstiitten side plates and acicular ferrite. Allotrimorphic ferrite and

WidmanstUtten side plates are generally not observed in high strength steel weld

metals due to the high alloying contcnt[ 18, 19].

Acicular ferrite is generally found to be positive for impact toughncss. It is called

acicular uue to its needle shape in two dimensional sections. In three dimensions,

it is founu to be in the form of lenticular plates with dimensions in the order of 5

to 10 pm in length and around I~lm in diameter. It nucleates on inclusions inside

the austenite grains. In literature it is sometimes described as an intragranlliarly

nllcleatcu bainite.

47

2.3.7 Bainite

Bainite was first discovered in the early 1930'5 by Davenport and Bain who were

investigating the isothermal transformations of austenite. It was called "Bainite"

by the staff at the United States Steel Corporation Laboratory after E.C. Bain who

had initiated the studies. From the early days it was known that there were two

types of bainite; upper bainite formed at high temperatures and lower bainite

formed at low temperatures[ 17]. A schematic of both, Upper and Lower bainite is

shown in Figure 2.4.

Carbon dilfusion Into austenite

Carbon sUp8rSalUtated plate

«ltMtl1ra /' ~ Catbon diflusion into

/' ~ austenite and carbide

~~';::::~,~ji'" precipitation in Ie-nite

..... . :'q;;;;/; /t?

',' ... ,' .... , ..... .

! Carbide preclpllalion from auslenite

LOWER BAINITE uppm BAINITE tHigh Temperalure) I Low Temper.lure)

Figure 2.4 Schematic of the transition path to both upper and lower bainite.

At transformation temperatures in the region of, or less than 690°C, carbon

docsn't have time to diffuse into austenite as ferrite nucleates on the austenite

48

grain boundaries. To cope with this, the carbon concentrates and redistributes at

the phase boundary since the solubility of carbon in bainitic ferrite is low « 0.02

wt.'70 [14]). As the carbon content increases at the phase boundary, it may reach a

high enough concentration at which cementite is able to nucleate and grow.

Depending on the level of carbon, cementite may be in the form discrete particles

(at low concentrations), or as continuous layers separating the ferrite plates (at

high concentrations).

Cementite or IvbC is an iron rich carbide where sometimes the iron atoms can be

replaced and other elements can he taken into solution. It has an orthorhomhic

structure with space group Pbnm. and Fe3C has lattice parameters a = 0.4523 nm,

b = 0.5088 nm and c = 0.6743 nm. Strong carbide forming elements such as Cr,

Mo. W, V and weaker clements such as Mn, Ni and Co can be found in solution

in M,C. It is formed by eg. eutcctoid and bainitic decomposition of austenite or

precipitation in ferrite during tcmpering.[ 13]

As each plate of bainitic ferrite grows it changes shape and is accompanied by a

large strain, which can be described as an invariant - plane strain with a large

shear component similar to that of martensite. The change of shape and

mechanism of growth indicate that the transformation mechanism is displacive

rather than reconstructive[21]. The two methods of atomic rearrangement,

reconstructive transformation and displacive transformation arc described in

literature. Reconstructive transformation is the breaking of all bonds and

rearranging the atoms into another alternative pattern while displacive

transformation is described as homogeneously deforming the original pattern into

a new crystal stmcture.

Displacive transformation also alters the macroscopic shape unless it is

constrained. During constrained rearrangement, transformation occurs by the

combination of elastic and plastic strains in the surrounding matrix. The product

49

grows in the form of thin plates, which reduces the strains. To conclude, bainite

fulfils all criteria for the displacive mechanism of transformation.

2.3.8 Martensite

At the fastest cooling rates, the austenite may reach the region of 400 DC or below

depending on alloying, before transformation. Transformation at these

temperatures is a diffusionless process producing lath martensite. No single atom

moves more than one lattice spacing. Martensite may be considered as a

supersaturated solution of carbon in a iron. It has a body centered tetragonal

(B.C.T.) structure. Since the transformation is from F.c.c. to B.C.C. and is also

diffusionless, the interstitial carbon atoms are limited to the one common

octahedral site in both the F.c.c. and B.C.C. structures[19]. The z-sites in this is

the octahedral sites common to both structures and will therefore be occupied by

carbon. As a result an expansion in the z direction is experienced. This is a

unidirectional distortion and bring about the change from B.C.C. to B.C.T. The

change to B.C.T. distinguishes martensite from supersaturated ferrite. In

supersaturated ferrite, all interstitial sites have equal probability of occupation.

When only substitutional solutes are present the transformed product is cubic

while it is generally tetragonal if there are interstitial solutes present. It was

reported that martensite in iron-carbon alloys with only -0.05 % carbon is

tetragonal, provided that the Ms temperature is sufficiently low to prevent

autotempering during the quench[18].

2.3.9 Effects of Tempering

Tempering is a term that comes originally from the heat treatment of martensite. It

generally describes how both the microstructure and properties change as a result

of holding a sample isothermally below the austenite formation temperature[ 16].

In multi-pass welding, the microstructure is deemed to be tempered once the weld

50

metal has been heated to temperatures below austenite formation as a result of the

deposition of a weld bead(s) on top of an underlying bead(s).

(I) Tempering of Martensite

qO{-CO"L-

Mainly there are four stages in the tempering of martensite. During the first stage

which takes place at low temperatures (up to 250°C), carbon in solid solution

begins to segregate to defects and forms clusters. It then precipitates to form

cementite or £-carbide which has a close packed hexagonal structure. Stage 2

happens in the region (230-300°C) where the majority of the carbon is

precipitated, £- carbide transforms to cementite which is more stable and any

austenite retained during quenching decomposes[14]. During Stage 3 cementite

appears and this can happen between approximately 100-300 dc. Also the

tetragonality in high carbon martensite disappears and the microstructure becomes

ferrite not supersaturated in carbon. Tempering at even higher temperatures (300-

700°C) results in Stage 4 where the cementite particles start coarsening, they lose

their crystallographic morphology and become spheroidised. Additional effects at

high temperatures are the recrystallisation of ferrite plates into equiaxed grains

which also promotes a reduction in the dislocation density[ 17]. Alloy carbides

begin to form in the range 500-600°C and their formation depends largely on the

alloying content and the presence of carbide forming elements.

The observed effects of tempering martensite on mechanical properties are

dependent on the carbon content within the martensite, the alloying content and

the tempering temperature[ 17]. For hardness the general trend is softening as the

tempering temperature is raised. A reduction in yield and tensile strength is also

normally observed while toughness is found to increase as tempering temperature

increases.

51

(II) Tempering of Bainite

There are important differences in the tempering of bainite from that of martensite

because bainite only contains a slight excess rather than a super saturation of

carbon in solution and it also auto tempers during transformation. A lot of the

carbon partitions out or precipitates during the bainite reaction[8, 12]. The carbon

is generally located within the cementite that forms which is usually coarser than

that formed from tempering martensite. Bainite is normally less sensitive to

tempering in comparison to martensite.

If bainite formation takes place at high temperatures, it may expenence some

recovery of the dislocation substmcture. Normally, only minor changes in

recovery, morphology or carbide precipitation occur during tempering. The

largest effect of tempering is observed when the plate-like stmcture of ferrite

changes to equiaxed ferrite. Spheroidisation and coarsening of cementite also take

place with this change[24J. Secondary hardening may also be experienced with

bainite at high temperatures but it is normally much more sluggish since the

cementite tends to be coarser in bainite.

Strength decreases with the tempering of bainite but the extent is not as dramatic

as with martensite because there is little or no carbon in solid solution. Carbon is

normally present in the form of carbide which contribute little to strength.

Changes in strength may be observed as the microstmcture coarsens or when

recrystalisation takes place at high temperatures and equiaxed grains of ferrite

replace the bainite plates. Minor changes may also occur due to the coarsening of

cementite particles or the recovery of the dislocation substmcture.

52

2.4 Alloying Element Addition

A typical analysis of 2.25Cr-IMo-0.25V steel is shown in Table2.3, as per

ASME compositional tolerances. The specific rolls of major alloying addition are

detailed below.

Table 2.3: Composition of2.25Cr-lMo-O.25V alloy steel (wt%,)

C Mn Si Cr Mo P S V Cb

SAW 0.05- 0.50- 0.05- 2.00- 0.90- 0.015Max 0.015Max 0.20- 0.010-

0.15 1.30 0.35 2.60 1.20 0.40 0.040

Solid solutions are formed when alloying clements such as carbon, manganese

are dissolved atom by atom either interstitially or substitutionally. When the

matrix is saturated with a particular element, a separate phase can form. As a weld

metal is either cooled or heated certain elements may inhibit phase transformation

in the steel[15.281. For example, austenite stabilisers are nickel, cobalt, carbon,

manganese and nitrogen while ferrite stabilisers are silicon, chromium, tungsten,

molybdenum, vanadium and niobium. Alloying also plays a role in mechanical

behaviour of the weld metal. Grain refinement is the most favoured strengthening

mechanism in high strength steel weld metals as it also contributes to toughness.

Precipitation strengthening is also a favoured strengthening mechanism[22,281. In

order to provide some resistance to brittle cleavage fracture, alloying with some

substitutional solutes helps to improve toughness. The manganese has very high

impact in improvement on toughness. The manganese improves micro structural

phases and mechanical properties.

53

2.4.1 Carbon

Carbon content is important to the overall strength and hardness of the weld

metal. The location of carbon atoms, whether they remain in solution or if they

are precipitated, determines whether the steel is martensitic or ferritic. The level

of carbon IS critical for optimizing microstmcture and mechanical

propcrties[22,28].

Greater additions of carbon lowers the martensite transformation temperature.

Increasing the amount of martensite with higher carbon in the HAZ raises the risk

of hydrogen cracking and also decrea~es toughness. At too low carbon contents

ferrite may form so carbon is made up for by increasing the alloying content.

The carbon content is also directly related to the risk of solidification cracking. At

high carbon levels greater amounts of Mn and lower levels of S arc required in

order to avoid this type of cracking. The carbon and sulphur contcnt is generally

kept low in welding consumables and solidification cracking is not a major

problem.

2.4.2 Manganese

The alloying content of manganese is very important in the solidification process

in high strength steel weld metals[16,28]. Large additions of these elements can

prevent the formation of 8-ferrite entirely and instead the weld mctal solidifies

directly to austenite.

Additional effects of manganese is that it gives strengthening through solid

solution hardening and grain refinement by lowering the austenite to ferrite

transformation temperature' Grain refinement also leads to increased toughness.

UnfOI1unately manganese tends to segregate which promotes variations in

54

hardness and microstructure[ 15,28]. It also forms inclusions that may be

detrimental to the toughness. Manganese is generally limited to minimize

solidification segregation. Manganese combines with sulphur and reduces the risk

of solidification cracking. Mn forming oxides and creating more number of

nucleation sites during solidification. Manganese increases high temperature

strength by solid solution hardening as well as toughness by grain refinement due

to favourable formation of suitable oxides and sulphides during solidification.

2.4.3 Chromium and Molybdenum

These elements are well known in the alloying of steel. Chromium stabilises

ferrite but slows down transformation rate. It also increases hardness and strength

and has a greater influence when manganese is at low concentrations. With

chromium additions, toughness falls as reported in studies of mechanical

properties of high strength steel weld metals[ 17,28]. It provides solid solution

strengthening and promotes carbide formation. Both chromium also increases

hardenability and gives both oxidation and corrosion resistance. Chromium and

molybdenum increase resistance to high temperature corrosion and are well

known for gi ving resistance against creep in heat resistant steels.

2.4.4 Silicon

Silicon contributes a large solid solution hardening factor. The addition of silicon

(in high concentrations, 1.5 wI. %) is reported to eliminate poor toughness in

bainitic steels. Silicon has poor solubility in cementite and retards its

precipitation.

2.4.5 Titanium, Niobium and Vanadium

Titanium, niobium and vanadium are microalloying elements. These elements can

form precipitates in the form of carbonitrides but are in very low quantities in

high strength steel weld metals[ 17,22]. If precipitation takes place, it contributes

55

to strength through precipitation hardening. Titanium and boron Increase

toughness at low temperatures. Titanium is also know to form oxides.

2.4.6 Oxygen, Nitrogen and H)'drogen

Oxygen, nitrogen and hydrogen are the three main absorbed gases within weld

metals. Oxygen reduces both the strength and toughness. It is difficult to

determine how much is in solid solution since it tends to be present in finely

dispersed oxygen rich microscopic inclusions. These microscopic inclusions are

an inevitable part of welding and in SAW their size is determined by the flux type

and energy input. Oxides can be in the form of iron-manganese oxide in neutral

and acidic weld metal with the latter containing some iron-manganese-silicate

oxides while the basic type weld metals almost only contain iron-manganese­

silicate oxides[28]. Oxygen also combines with iron to form FeO and separates

out as slag.

Nitrogen raises the strength but reduces the toughness. Nitrogen is normally

present in ppm concentration but may still have a large effect on the mechanical

properties. Too long an arc may lead to excess nitrogen pick up. Excess nitrogen

is ejected in the form of gas or may cause porosity.

Hydrogen has a high solubility in iron because of its small atomic radius and

dissolves easily in the solidifying weld metal. Although it is present in solution

while welding it diffuses out as the weld metal cools[24,28]. When practical any

remaining hydrogen can be removed by a low temperature « 300) soaking heat

treatment. Hydrogen within the weld metal reduces ductility.

56

;fv ,

_J .:::.'. :-

2.5 Effect of Carbides on Microstructure

The alloy content of this grade of steel exceeds the soluhility limit of the

BCC matrix phase and therefore 2.2SCr-IMo-O.2SV is a multi-phase material

under equilibrium conditions. The excess alloy contact exists in the form of

carbides distributed throughout the ferrite matrix. The alloy carbides can dissolve

significant amounts of other alloying elements to form non-stoichiometric alloy

carbide phases[26,28].

When heated to service temperatures, the carbides will change composition and

morphology to approach equilibrium within the service life The stable phase

eventually reached is dependent on several factors, such as initial structure and

compositional segregation are of the M"C or M2}C6 type where the M symbolises

the sum of all metallic species in the carhide[2S,36].

The progression towers the stable carhide structure involves several steps trough

intermediate carbide phase. Early work by Baker and Nutting proposed the

following sequence of carbide precipitation, by observing the tempering response

in both normalized and quenched material[2S,36]:

More recent experimentation by Tsai and Yang agreed with this general sequence

hut failed to observe M6C and named the final product of the carbide progression

as M2}C6. in contrast, Williams describes M6C as intermediate to the towards

structures with a higher metal content. The properties and precipitation

mechanisms of these carbides are summarized below[26,28,36]:

M3C: This carbide is initially formed as Fe]C, and can dissolve other alloy

constitutes such as Cr to replace up to around 20% of the initial Fe. Mo can also

be accommodated in the carbide structure hut only up to approximately 4%. Apart

57

from €-carbide, M)C is the least thermodynamically stable of the carbides, it is

quickly replaced by higher alloy carbides during heating.

M 2C: This carbide is based on M02C and can dissolve significant amounts of

Cr(up to 30% of its weight). It can be precipitated from both the quenched and

normalized structure. This carbide is founded to precipitate by intergranular

nucleation directly from the ferrite[26,36]. Initially, M02C is trough to be coherent

with the matrix, before developing into needle-shaped, non-coherent precipitates.

M)C): This carbide is based on the Cr)C) structure is capable of dissolving

significant amounts of Fe. The precipitation of this carbide occurs in the vicinity

of the existing Fe)C, nucleating at ferrite/Fc)C interfaces[26,36]. Through studies

of Cr-C alloys, the Cr)C) carbide was shown to precipitate and coarsen much

faster than other alloy carbides and as a result is not expected to produces

signitlcant secondary hardening.

M 2.,C6 : This carbide is through to be based on either Cr2JC6 or Fe21 M02C6• The

precipitation of this phase occurs at expense of existing Fe)C and Mo2e. Baker

and Nutting showed it was only developed in quenched in quenched structures or

bainitic areas of this alloy[25,36]. It does not require the dissolution of Cr)C.l and

is not observed in areas of existing Cr)C3 . This carbide nucleates and grows as

plate-like particles.

M 6C: This carbide has been described as FeJMo,C by Qu and Kuo and as varying

between Fe.lMo)C and Fe.Mo1C by Olsen and north. The structure is also c,lpable

of dissolving significant amounts of Cr[25]. This is the most alloyed of the

carbide reaction path outlined above by several studies. It form as large faceted

particles at grain boundaries. Baker and Nutting showed this carbide to form in

58

both the bainite and ferrite of a normalized structure. MoC is trough to provide

less dispersion strengthening than other carbides.

Research by Williams showed that the M23C6, M6C and a M7C3 phases can all

exist simultaneously in the same alloy sample. Different grains can progress along

the carbide evolution path at different rates due to small-scale heterogeneity.

The rate of degradation and the morphology of the precipitated and reprecipitated

carbides show a complex response to temperature and initial microstructure. Since

the carbide development is temperature and initial microstructure. Since the

carbide development is dependant on the initial microstructure, a heterogeneous

microstructure such as a weldment will exhibit a variety of carbide distribution

throughout its different zones[26,36].

In sLlIllmary, the dispersed carbide phases in a 2.2SCr-Mo-O.2SV alloy are

metastable in both the as-welded and normalized(parent metal) conditions, and

will change progressively towards more stable form under thermal activation.

2.6 Factors Affecting Toughness of Weld

There are various factors affecting the toughness of weld. Some of these are

discussed below.

2.6.1. Grain size and welding technique

Grain size and percentage of fine grained microstructure are very important due to

the fact that primary grains after solidification are partially refined by each of

sLiccessive pass during welding [27]. The extent of this refinement influences the

weld properties. Thus, geometry of each bead has to be considered as a decisive

factor of intluence since the extent of refineIllent depends directly on thickness of

each bead or width / thickness ratio. The thinner the weld passes, larger are the

59

areas, which are refined by each successive weld pass and better are the toughness

values. For this reason, it is very important to adhere closely to proper welding

procedure.

Also, specific welding technique will affect the toughness when dealing with

vertical uphill (3G) or other restrictive welding positions.

2.6.2. Bead deposition layout

It is also important from toughness point of view. It is observed that the weld

metal specimens with the notch placed at the centre of the overlap zone (refined

zone) showed in the most cases much higher toughness than those with the notch

in the centre of the bead.

2.6.3. lIeat input

Heat input should be limited to provide higher toughness of weld and better

control of delayed cracking. 1I0wever, it is believed that welding parameter

themselves, rather than net heat input are responsible for producing a

microstructure of acceptable toughness[27J.

2.6.4. Type of current

One of the major factors that affects the toughness of weld metal is concerned

with type of current. Welding with alternating current gives better toughness of

weld than that with direct current due to the fact that in the fonner one, weld

contains lower oxygen and better agitation of weld metal..

2.6.5. Chemical composition of Filler Metal

Welding wire chemical composition is usually determined by the base metal and

by service requirements. For higher toughness, welding wire should provide a

very low oxygen and nitrogen in the weld metal. Also, impurities such as P, As,

Sb, Sn etc, in welding wire should be limited to a very low value[25J.

60

There is an improvement in toughness of weld by increasing Mn content. It is due

to the fact that increasing the content of this elements leads to the microstmcture

relinement because of lowering the BS (start of bainitic transformation)

temperature.

2.6.6. Flux Passivity

Flux with high basicity in agglomerated or bonded form is good for SAW since it

gives lower number of inclusions and thereby better toughness of weld metal.

Also, basic fluxes reduces silicon transfer from flux to weld metal and provide

weld metal with an oxygen content. It should also provide very low Sulfur ,

Phosphoms and residual elements transferred to weld metal.

2.6.7 Welding Metallurgy

To establish the close relationship between weld nllcro stmcture and impact

toughness of weld, following important points are strictly observed during

fabrication of Cr-Mo-V pressure vessels steel weld for hydro cracker reactor[26]

I) Base metal composition.

2) Composition of consumables.

3) Welding process and parameters.

4) Deposition layout.

5) Pre and post weld heat treatment.

The above description provides necessary insight for interdependent requirement

of fabrication process and metallurgical aspects of hydrocracker equipment.

61

REFERENCES:

ahavnaga r ! I .. ;"crsity Li':

BHA" NAGAR.

I. API, "Steels for Hydrogen Services at elevated Temperatures and Pressure in Petroleum Refineries and Petrochemical Plants," API Pub!. 941 (1983)

2. R.D.Stout, " Weldability of Steel ", Welding Research Council, FouI1h

edition.

3. R.C.Gupta, "Welding Engineering Science & Metallurgy"

4. N.K.Roy, "Cr-Mo Alloy Steels - Their Elevated Temperature Properties".

5. ASME Sec IX, "Welding And Brazing Qualification"

6. N.K. Roy., Gramya research analysis institute. "Procedure Specification for Metal Arc Welding of Low Alloy Steels."

7. S.V.Nadkarni, Advani-Orlikon Limited, " Modern Arc Welding Technology".

8. Dr.A.K.Lahiri, Kribhco Ltd., Hazira, Surat, Gujarat, "Material Selection And Performance in Fertilizer Industries".

9. WRC Bulletin." 2.25Cr-1 Mo Steel For Vessel"

10. M.W.Hardie, "Aspects Of Welding 2.25cr-Imo Steel For Pressure Vessel"

II. ASM - Welding handbook., "Weldability of Steels"

12. Vijendra Singh, " Physical Metallurgy"

13. J.F.Henry, "Cr-Mo Welds - Science Failure" Lesson4, AlstOIll, Jan.2oo6

14. J.F.Henry, "The Metallurgy of the creep strength enhance ferritic steels ",

Lesson 5, Alstom , Jan.2006

IS. O.P.Khanna, "Welding Technology".

16. Welding Inspector Course - By I!W, 26Ih_30'h June, 2002

17. R.KJain, "Production Technology"

18. Dr. R.S.Parmar, "Welding engineering and Tehnology"

62

----------- --

19. Jefferson's "Welding Encyclopedia " Eighteen Edition, Edition by -

Robert L.O'Brien, AWS, 1997

20. M.Khandapetov, "Welding And CUlling of Metals"

21. O.P.Khanna , " Material Science and Metallurgy"

22. G.B.S.Narang, " Material Science"

23. Lancaster, " Physical Metallurgy"

24. R.PJoeypal, " Welding Fabrication and Inspection Techniques Industry

based" By Welding Research Council, on 11th - 25th June, 1984

25. Vitek, J.M. and klueh, R.I-I., "Precipitation reaction during the heat treatment of ferritic steels," mertal!. Trans. A, 14a( 6), 1047-1055(1983).

26. Wood head, J .1-1. and Quarrell, A.G., 'The role of carbide in low alloy creep resisting steels," climax molybdenum co. ltd., London, 1965.

27. R.Ravi & S. DUlla " Selection of electrodes for welding of Stainless Stcel

", D&H Sec heron

28. J.T.McGrath, J.A.Gianello, and R.f.Orr "Welding Procedures for the

Fabrication of Thick Section 2.25 Cr- 1 Mo Pressure Vessel Steel"

29. Journel of ISNT" NDT & Evaluation "- Vol 6, Issue I, Scpo 2007

30. Journel of ISNT" NDT & Evaluation "- Vol 6, Issuc 2, Sep. 2007

31. Y.K.Park, K.S. Kim, Y.K.Chung, J.J.Park, " Creep Crack Growth In 20Cr­Mo-V 12 I Steel And Its Weld Joint".

32. lun-ichi Shimomura, Yoshifumi Nakado, Syozaburo Nakado, syuzo Ueda "IMPROVEMENT IN RESISTANCE TO DIS BONDING OF STAINLESS STEEL-OVERLAYED 2 ,;.\ CR-I MO STEELS"

33. BOJAN A.MARINKOVIK, ROBERTO R. DE A VILLEZ, SIMONE KESSLER BARROS, FERNANDO C. RIZZO ASSUNCAO, 'THERMODYNAMIC VALUATION OF CARBIDE PRECIPITATES IN 2.25CR-1.0 MO STEEL FOR DETERMINATION OF SERVICE DEGRADA nON"

63

34. A.M.Adbel-Latif, J.M.Corbett, D.Sydeny and D.M.R. Taplin, "Effects of Microstructural Degradation on Creep life Prediction of 2.25Cr-IMo Steel", Advance in fracture research institute (Fracture 81), Proceedings of the Fifth International Conference On Fracture Research, Cannes, France, march 29-April 3, vol.l, pp.1613-1620.

35. M.DeWitte and C.Coussement, "Behavior of 2.25Cr-IMo Base Material and (Repair) Weldments at High Temperature", Proceedings of the fifth International Conference on Creep of Materials, Lake Buena Vista, Florida, U.S.A., May 18-21, pp. 417-422.

36. B.E. Peddle and c.A. Pickles , "Carbide Development In the Heat Affected Zone of Tempered and Post weld heat treated 2.25Cr-1 Mo Steel Weldments"

37. J.G.NAWROCKI, J.N.DUPONT,C.V.ROBINO, AND A.R.MARDER "THE STRESS RELIEF CRACKING SUSCEPTIBILITY OF A NEW FERRITIC STEEL - PART 2: MULTIPLE PASS HEAT AFFECTED ZONE SIMULATIONS"

38. ASM handbook vol. 6, handbook of joining and brazing, ASM international materials park, Ohio.

39. Shiro Nose, Masato Yamada, Toshiaki Fukada, and Tadamichi Sakai Energy and Chemical Plant Group, Kobe Steel, Ltd., Takasago, Hyogo, Japan & E.H. Nicolls, A.1. Bagdasarian, and A.G. Imgram, Chevron Research & Technulogy Co., Richmond, CA, U.S.A. "FABRICATION OF A HYDROPROCESSING REACTOR APPLYING 2.25 CR-I MO-V­CB-CA STEEL"

40. Parvathavarthini, N.; Saroja, S.; Dayal, R.K.; Khatak, H.S. j. Nucl. Mater.; v. 288, 187-196,2001.

41. Narcosis. J.G., Dupont, J.N., Robino, C.V., and marder, A.R. 2000. "The stress relief cracking susceptibility of a new ferritic steel part - I : single pass heat affected zone simulations.

42. Kessler, S. Msc Dissertation, DCMM, PUC-Rio de Janerio, 2000. (in Portuguese)

43. Sundman, B. Thermo-Calc Version L- User's Guide, 1998.

64

44. Hillert, M. Phase Equilibrium, Phase Diagrams and Phase Transformation: A Thermodynamic Basics, Cambridge University Press, 1998.

45. Nose S, Yam ado T, Yamada M., Sakai T, Normura T, and Aoyama T, 1998, "Highly Reliance Time of Flight Diffraction UT Technique and its Application to Pressue Vessel Fabrication", ASMEIPVP.

46. Lundin C.D., and Khan, K. K., 1996, "Fundamental Studies of the Metallurgical Causes and Mitigation of reheat cracking in I 1,4 Cr-I/2Mo and 2 1/4Cr-IMo Steel WRC Bulletin 409, WRC.

47. Honeycombe, R.W.K., and Bhadeshia, H.K.DH. 1996. STEELS: microstmctures and properties, 2nd ed., Halstead Press, New York, N.Y.

48. 1.D. DobIs, G.M.Buchiem, D.A.Osage."Material Selection & Design", MP/dec-1995. page 61-64.

49. Yamada M., Sakai T., and Nose S., 1995, "Development of Cr-mo-v-Cb­ca Steel for high Pressure a: I High Temperature Hydrogenation Reactors", PVP-Vol. 315, SP 385-388, ASME.

50. Moss, C.1.; Kelly P.M. Fatigue and Fracture Engineering Materials and Structures, v. 17, n.3, p.369-380,1994.

51. Thomson, R.c.; Bhadeshia, H.K.D.H. Material Science and Technology, v. 10, p. 193-203, 1994.

52. J.Janovec, A.Vyrostkova and M.Svoboda, "Inl1uence of Tempering Temperature on the stability of carbides phase in Metallurgical and Materials Transaction, 1994, vol. 25A, pp. 267-275.

53. Sakai t, Yamada M, Akasaka K, Pukada T, and Nose s, 1994, "Application of low temperature PWHT to advanced Cr-Mo steel", MPC 2ND International conference on interaction of steels with hydrogen in petroleum industry pressure vessel and pipeline service, pp 79-87, MPC.

54. Yang, 1.R.; Huang, c.Y.; I'luang, C.N.; Horng, 1.L. Material Characterization., v. 30, p. 75-88, 1993.

55. Fukada t. Nose S., Yamada M., and Terada, 1992- "Submerged arc welding of thick 2 1,4 Cr-I Mo steel using ultra low hydrogen type bonded l1ux" presented on MPC.

65

56. e.F. Etienne, J.H.Hcerings, J.e. van Wortel, L.B. Dufour, G.M. van Dijk and J .H.Fokkens, "Dutch Developments in Design and Behavior of Creep Loaded Structures", Proceeding of the fifth International Conference on Creep of Materials, Lake Buena Vista, Florida, U.S.A. May 18-21, 1992, pp.453-459.

57. e.D.Buscemi, B.LJack, J.W.Skogsberg and WE Erwin, ''Temper Embrittlement in 2.25Cr-IMo Steels after 75,OOO-Hour Isothermal Aging", Journal of Engineering Materials and Technology, July 1991, Vol. 113, pp. 329-335

58. T.W.Lau and D.Hartwick, "Waving Welding Techniques: Their Effect On The Microstructure And Hardness Of 2.25Cr-IMo steels", Ontario Hydro Research Division Report No. M91-61-K, 1991.

59. T.W.Lau and D.Hartwick, "The Effects of Thermal Cycles on the Hardness and the Microstructures of Repair Weldments of a 2.25Cr-1 Mo Steel", Ontario Hydro Research Division Report No. M90-33-K, 1990.

60. Lundin, e.D.; Khan, KK.; Yang. D.; Hilton, S.; Zielke, W. WRC Bulletin, 354,1990.

61. Imgram, A.G., Ibarra S., and Prager M.1990, "A vanadium modified 2 1,4

Cr-I Mo steel with superior performance in creep and hydrogen service", New Alloys for Pressure Vessels and Piping. PVP Volume 201, MPC Volume 31, ASME, #H00620.

62. Katsuma M, 1989 "Study on embrittlement of large scale forgings and its prevention", Doctor thesis Konkorn Univ.

63. J. Shimomura, E. Sugie, Y. Nakano and S. Ueda: Aturyoku Gijitsu, 27 (1989),22

64. M.e.Coleman, "High Temperature Behavior of Ferritic Pipe welds; Experience of Long-Term Testing", International Journal of Pressure Vessels and piping, 1989. vol. 39, pp. \09-118.

65. SJ.Brett, "The Weld repair of Cr-Mo-V Turbine Casings using high nickel electrodes without heat treatment", Materials Development in Turbo-Machinery Design, Second Parson International Turbine Conference, The institute of Metals, London, 1989, pp. 166-172.

66

66. J. Shimomura, Y. Nakano, S. Nakano, J. Tateishi and S Ueda: Proc, 6th

Int. Conf. On Pressure vessels Technology, Beijing, Pergamon Press, (1988), 10 13.

67. G.S.Kim, J.E. Indacochea and T.D. Spry, "Weldability Studies in Cr-Mo­V Turbine Rotor Steel", Journal of Materials Engineering,1988, Vol.lO, pp. 117-132

68. J. Shimomura, M. Miki, Y. Nakano and S. Umeda: Proc. 5TH Japan­Czechoslovakia Joint Symp., ISIJ, Tokyo, (1987),183.

69. Sakai T., Asami K., Kondo, N., and Hayashi, 1987, "Effects of carbide forming elements on hydrogen crack and embrittlement in 2.25 Cr-I Mo steels", ibid, Vol.73 Sp.372-379

70. Furusawa, J.; Watanabe, S. The Sumitomo Search, v. 35,1987.

71. Y.J.Kim and J.W. Prince, "Temper Bead Weld HAZ Properties in A516-70 Steels", Journal of Engineering Materials and Technology, 1987, vol. 109, pp. 157-163.

72. T. Sakai, K. Asami. W. Kondoh and F. Hayashi: Tetsuto-Hagane, 72 (1987),372.

73. T. Ishiguro, K. Ohnishi and J. Watanabe: Tetsu-to-Hagard, 72 (1986),70.

74. Senior, B.A. Nobble, F.W., and Eyre, B. L. 1986. Acta metallurgica 34(7): 1321-1327.

75. Sakai T., Takagi I., and Asami K., 1986, "Effect of carbides and inclusions on internal hydrogen embrittlement of Cr-Mo pressure vessel steel", Journal of the Iron and Steel Institute of Japan, Vol.72, pp. 1375-1382.

76. T. Imanaka, J. Shimomura, S. Nakano and k. yasuda Kawasaki steel tech. Rep., (1985) No. 13, 109. T. Ishiguro, K. Ohnishi and J. Watanabe: Tetsu­to-Hagard 71 (1985) 986.

77. T. Ishiguro, K. Ohnishi and J. Watanabe: Tetsu-to-Hagard 71(1985) 986.

78. T.lshiguro and K. Ohnishi: Tetsu-To-Hagard, 71(1985) S1304.

67

92. Sakai T, Asami K, ct aI., 1982, "Hydrogen induced disbonding of weld overlay in pressure vessel and its prevention", Current Solutions to Hydrogen Problems in Steels, pp 340, ASM.

93. Murakami, Y.; Nomura, T.; Watanabe, J. Application of2.25 Cr-I Mo Steel for Thick wall Pressure Vessel, ASTM STP 755, American Society for Testing Materials, p. 383-417, 1982.

94. C.A.Hippsley, "Precipitation Sequences in lIeat Affected Zone of 2.25Cr­IMo Steel During Stress Relief Heat Treatment", Metal science, 1981, vol. IS, pp. 137-147.

95. R.L.Klueh, "Creep-Rupture Behavior of a Bainitic 2.25cr-1 mo steel", International Journal of Pressure Vessel And Piping, 1980, vol.8, pp.165-185.

96. Inuoe, A.; Masumoto, T. Metallurgical Transaction A, v. II, p. 739-747, 1980.

97. K. Naito, H. Okada, J. Watanabe and K. Ohnishi: Aturyoku Gijitsu, 18 (1980),47.

98. R.F.Hoare, "Welding Procedure and Evaluation of Heavy Section 2.2SCr­I Mo Steel weldments", Canada Center of Mineral and Energy Technology, Physical Metallurgy Research Laboratories (Canada), ERP/PMRL 81-75(TR), 1981

99. Kihara, S., Newkirk, J.B. Ohtomo, A., and saiga, Y. 1980. Metallurgical transition A. ,IIA(6): 1019-1031.

100. Curry.D.A., and Pratt, P.L. 1979, material science and engineering. Vol. 37, pp. 223-235.

101. J.Orr, F.R. Beckitt and G.D. Fawkes, 'The Physical Metallurgy of Chromium-Molybdenum Steels for Fast Reactor Boilers", Ferritic Steels for Fast reactor steam Generators, BNES, London. 1978, pp. 91-109

102. Hillert, M. The uses of Gibbs Free Energy - Composition Diagrams, in the Lectures on the Theory of Phase Transformations, Ed. H.T. Aaronson, The Metals Society of AIME, 1975.

103. Meitzner, c.F. 1975. WRC Bulletin 211, pp. 1-17

104. Vinckier, A.G., and pense, A.W., 1974. WRC Bulletin 197.

69

105. R.G.Baker and J.Nutting, "The Tempering Of 2.25Cr-IMo Steel After Quenching And Normalizing", Journal Of The Iron And Steel Institute, 1972, vol. 210, pp. 257-268.

106. K.W.Andrews, H.Hughes, and DJ.Dyson, "Constitution Diagrams For Cr-Mo-V steels", Journal of the Iron and Steel Institute, 1972, vo1.21O, pp.337-350.

107. M.C.Murphy and G.D.Branch, "Metallurgical Change in 2.25Cr-1 Mo steels during creep rupture test", journal of the iron and steel institute, 1971, vol. 209, pp. 546-561

108. K.W.Andrews, DJ.Dyson and S.R.Keown. "Interpretation of electron diffraction patterns", sccond edition, plenum press, 1971, pp. 180-203.

109. Lundin, C.D., Liu, P., Qiao, c.Y.P., Zhou, G., Khan, K.K., and Prager, M. 1966. WRC bulletin 411.

110. L.M.T. Hopkin, D.Murray and R.Duval, "Creep Properties of Some Butt Welds in Steam pipes", Journal of the Iron and Steel Institute, 1965, vol. 203, pp. 819-825.

III. G. Lange and W. Hofman: Arch. Eisenhuttenwes., 37( 1966) 391.

112. J.lLWoodhead and A.G.Quarrell, "Role of carbides in low-alloy creep resisting steels", journal of the iron and stcel institute, 1965, vo1.203, pp. 605-620.

113. K.F.Hale, "An electron Microscope Study of changes in a 2.25% Cr, I %M Super-Heater Tube Steel During Tempering and Creep", proceedings of the fourth international conference on electron microscopy, springcr­verlag, berlinm 1960. yol.l \, pp. 650-658.

114. E.W. Johnson and M.L.Hill: Trans. Am. Inst. Mix. Metall. Pet. Engg .• 218(1960), 1104.

115. Baker.R.G .. and Nutting. J. 1959. journal of the iron and stcel institute. July. Pp. 257-268.

116. Nippes, E.F., and Nelson. E.C. 1958. Predictions of weld heat affected zone microstructures from conditions cooling transformation data. Welding journal 37(7): 289-s to 294-s.

70

1l7. Seal, A.K. and honcycomb, R.W.K., "Carbide preCIpItation in several steels containing in chromium and vanadium," 1. iron steel ins!., 9-15(JAN.1958).

118. Pctch, N.J. 1953. JIST. Vol. 173, p. 25.

119. Hall, E.O. 1951. Proc. Phys. Soc. B. Vol. 64. P. 747.

120. Nippes, E.F., Merrill, L.L., and savage, W.F. 1949. cooling rates in arc welds in 1/2 -in plates. Welding journal 218( II): 556-5 to 564-5

121. G. Phragmen: lernkont, Ann., 128(1944),537.

71