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5.

Welding of High-Alloy Steels,

Corrosion


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5. Welding of High-Alloy Steels, Corrosion 58

Basically stainless steels are characterised by a chromium content of at least 12%. Figure

5.1 shows a classification

of corrosion resistant

steels. They can be sin-

gled out as heat- and

scale-resistant and

stainless steels, depend-

ing on service tempera-

ture. Stainless steels are

used at room temperature

conditions and for water-

based media, whilst heat-

and scale-resistant steels

are applied in elevated

temperatures and gaseous

media.

Depending on their microstructure, the alloys can be divided into perlitic-martensitic, ferritic,

and austenitic steels. Perlitic-martensitic steels have a high strength and a high wear resis-

tance, they are used e.g. as knife steels. Ferritic and corrosion resistant steels are mainly

used as plates for household appliances and other decorative purposes.

The most important group are austenitic steels, which can be used for very many applications

and which are corrosion resistant against most media. They have a very high low tempera-

ture impact resistance.

Based on the simple Fe-C

phase diagram (left figure),

Figure 5.2 shows the ef-

fects of two different

groups of alloying elements

on the equilibrium diagram.

Ferrite developers with

chromium as the most im-

portant element cause astrong reduction of the aus-

Classification of Corrosion-Resistant Steels

non-stabilized

(austenite withdelta-ferrite)X12CrNi18-8

stabilized

(austenite withoutdelta-ferrite)

X8CrNiNb16-13

ferritic austenitic

stainlesssteels scale- and heat-resistantsteels

corrosion-resistant steels

semi-ferritic ferritic-austenitic

X40Cr13 X10Cr13 X8Cr13 X20CrNiSi25-4

perliticmartensitic

ISF 2002br-er-06-01e.cdr

Figure 5.1

Modifications to the Fe-C Diagramby Alloy Elements

ChromiumVanadiumMolybdenumAluminiumSilicon

NickelManganeseCobalt

Alloy elements in %Alloy elements in % Alloy elements in %

T

A4

A3

T

A4

A3

T

A4

A3


Figure 5.2


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tenite area, partly with downward equilibrium line according to Figure 5.2 (central figure).

With a certain content of the related element, there is a transformation-free, purely ferritic

steel.

An opposite effect provide austenite developers. In addition to carbon, the most typical mem-

ber of this group is nickel.

Austenite developers cause an extension of

the austenite area to Figure 5.2 (right figure)

and form a purely austenitic and transforma-

tion-free steel.

The table in Figure 5.3 summarises the ef-

fects of some selected elements on high alloy

steels.

The binary system Fe-Cr in Figure 5.4 shows

the influence of chromium on the iron lattice.

Starting with about 12% Cr, there is no more

transformation into the cubic face-centred

lattice, the steel solidifies purely as ferritic. In

the temperature range between 800 and

500C this system contains the intermetallic

-phase, which decomposes in the lower

temperature range into a low-chromium -

solid solution and a chromium-rich -solid

solution. Both, the development of the -phase and of the unary --decomposition cause a

Effects of Some Elementsin Cr-Ni Steel

Element Steel type, no. Effect

Carbon

l

l

l

All types

l

l

l

Increases the strength, supports development

of precipitants which reduce corrosion

resistance, increasing C content reduces

critical cooling rate

Chromium

l

All types

l

Works as ferrite developer, increases

oxidation- and corrosion-resistance

Nickel

l

l

All types Works as austenite developer, increases

toughness at low temperature, grain-refining

Oxygen

l Special types l

Works as strong austenite developer

(20 to 30 times stronger than Nickel)

Niobium

l

1.4511,1.4550,

1.4580 u.a.

Binds carbon and decreases tendency to

intergranular corrosion

Manganese

l

l

All types

l

l

Increases austenite stabilization, reduces hot

crack tendency by formation of manganese

sulphide

Molybdenum

l

l

1.4401,1.4404,

1.4435 and others.

l

Improves creep- and corrosion-resistance

against reducing media, acts as ferrite

developer

Phosphorus,

selenium, or

sulphur l

1.4005, 1.4104,

1.4305

l

l

Improve machinability, lower weldability,

reduce slightly corrosion resistance

Silicon l

l

All types

l

l

Improves scale resistance, acts as ferrite

developer, all types are alloyed with small

contents for desoxidation

Titanium

l

l

1.4510, 1.4541,

1.4571 and others

l

Binds carbon, decreases tendency to

intergranular corrosion, acts as a grain refiner

and as ferrite developer

Aluminium

l

Type 17-7 PH

l

Works as strong ferrite developer, mainly

used as heat ageing additive

Copperl

l

l

Type 17-7 PH,1.4505, 1.4506

l

l

Improves corrosion resistance against certain

media, decreases tendency to stress

corrosion cracking, improves ageing

ISF 2002br-er06-03e.cdr

Figure 5.3

Figure 5.4


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strong embrittlement. With higher alloy steels, the diffusion speed is greatly reduced, there-

fore both processes require a relatively long dwell time. In case of technical cooling, such

embrittlement processes are suppressed by an increased cooling speed.

Nickel is a strong austenite developer, see Figure 5.5 Nickel and iron develop in this system

under elevated temperature a complete series of face-centred cubic solid solutions. Also in

the binary system Fe-Ni

decomposition processes

in the lower temperature

range take place.

Along two cuts through the

ternary system Fe-Cr-Ni,

Figure 5.6 shows the most

important phases which

develop in high alloy steels.

A solidifying alloy with 20%

Cr and 10% Ni (left figure)

forms at first -ferrite. -

ferrite is, analogous to the

Fe-C diagram, the primary

from the melt solidifying

body-centred cubic solid

solution. However -ferrite

is developed by transfor-

mation of the austenite, but

is of the same structure

from the crystallographic

point of view, see Figure

5.4.

Binary System Fe - Ni

30Fe 10 20 40 50 60 70 80 90 Ni

0

200

400

600

800

1000

1200

1400

1600

Temp

erature

C

%

Nickel

Fe Ni3

FeNi 3

S+S+


Figure 5.5

Sections of the Ternary System Fe-Cr-Ni

700

800

900

1000

1100

1200

1300

1400

1500

1600

0 5 10 15 20 % Ni

% Cr30 25 20 15 10

70 % Fe

0 5 10 15 20 25700

800

900

1000

1100

1200

1300

1400

1500

1600

40 35 30 25 20 15

% Ni

% Cr

60 % Fe

Temperature

C C SS

S+

SSS+

S+

S+


Figure 5.6


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7/19


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The ferrite content can only be measured with a relatively large dispersal, therefore DeLong

proposed to base a measurement procedure on standardized specimens. Such a system

makes it possible to measure comparable values which don't have to match the real ferrite

content. Based on these measurement values, the ferrite content is no longer given in per-

centage, but steels are grouped by ferrite numbers. In addition to ferrite numbers, DeLong

proposed a reworked Schaeffler diagram where the ferrite number can be determined by the

chemical composition, Figure 5.13. Moreover, DeLong has considered the influence of nitro-

gen as a strong austenite developer (effects are comparable with influence of carbon). Later

on, nitrogen was included into the nickel-equivalent of the Schaeffler diagram.

The most important feature

of high alloy steels is their

corrosion resistance start-

ing with a Cr content of

12%. In addition to the

problems during welding

described by the Schaeffler

diagram, these steels can

be negatively affected with

view to their corrosion re-

sistance caused by the

welding process. Figure

5.14 shows schematically

the processes of electro-

lytic corrosion under a

drop of water on a piece of

iron. In such a system a

potential difference is a

precondition for the devel-

opment of a local element

consisting of an anode and

a cathode. To develop

De Long Diagram

16 17 18 19 20 21 22 23 24 25 26 27Chromium-equivalent = %Cr + %Mo + 1,5 x %Si + 0,5 x %Nb

Nickel-equivalent=%Ni+30x%C+30x%N+0,5x

%Mn

21

20

19

18

17

16

15

14

13

12

11

10

austenite

Schaeffler-austenite-martensite-line

austenite + ferrite

form

erly

mag

netic

allym

easu

red

ferrite

conten

tsin

vol.-%

ferritenu

mbe

r

2%

4%

6%

7,6%

9,2%

10,7%

12,3%

13,8%

0%

0

2

4

68

1012141618


Figure 5.13

Corrosion Under a Drop of Water

air

water

Fe(OH)3

iron

2Fe +O+H O 2Fe +2OH++ +++ -

2

H O2

O

OH-

cathode

anode

2Fe 2Fe +4e ++ -

4e-

O +2H O+4e 4OH2 2- -

O2 OH

Fe+++

2Fe++


Figure 5.14


9/19


10/19


If crevice corrosion is pre-

sent, corrosion products built

up in the root of the gap and

oxygen has no access to

restore the passive layer.

Thus narrow gaps where the

corrosive medium can ac-

cumulate are to be avoided

by introducing a suitable de-

sign, Figure 5.16.

With pitting corrosion, the

chemical composition of the

attacking medium causes a

local break-up of the passive layer. Especially salts, preferably Clions, show this behaviour.

This local attack causes a dissolution of the material on the damaged points, a depression

develops. Corrosion products accumulate in this depression, and the access of oxygen to the

bottom of the hole is obstructed. However, oxygen is required to develop the passive layer,

therefore this layer cannot be completely cured and pitting occurs, Figure 5.17.

Stress-corrosion cracking occurs when the material displaces under stress and the passive

layer tears, Figure 5.18. Now the unprotected area is subjected to corrosion, metal is dis-

solved and the passive

layer redevelops (figures 1-

3). The repeated displace-

ment and repassivation

causes a crack propaga-

tion. Stress corrosion

cracking takes mainly

place in chloride solutions.

The crack propagation is

transglobular, i.e. it does

not follow the grainboundaries.

Pitting Corrosion of a

Storage Container

Steel

br-er-06-17e.cdr

Figure 5.17

Model of Crack PropagationThrough Stress Corrosion Cracking

1 2 3 4 5 6

121110987

offset; passive layer; metal surface; dislocation

br-er-06-18e.cdr

Figure 5.18


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Figure 5.19 shows the expansion-rate dependence of stress corrosion cracking. With very

low expansion-rates, a curing of the passive layer is fast enough to arrest the crack. With

very high expansion-rates, the failure of the specimen originates from a ductile fracture. In

the intermediate range, the material damage is due to stress corrosion cracking.

Figure 5.20 shows an example of crack propagation at transglobular stress corrosion crack-

ing. A crack propagation speed is between 0,05 to 1 mm/h for steels with 18 - 20% Cr and 8 -

20% Ni. With view to welding it is important to know that already residual welding stresses

may release stress corrosion cracking.

The most important problem in the field of welding is intergranular corrosion (IC).

It is caused by precipitation of chromium carbides on grain boundaries.

Although a high solubility of carbon in the austenite can be expected, see Fe-C diagram, the

carbon content in high alloyed Cr-Ni steels is limited to approximately 0,02% at room tem-

perature, Figure 5.21.

TransgranularStress Corrosion Cracking


Figure 5.20

Influence of Elongation Speed onSensitivity to Stress Corrosion Cracking

SpRK

completecover layer tough fracture

Sensitivi

tyto

stresscorrosion

cracking

Elongation speed

2 1

T=RT

ISF 2002br-er06-19E.cdr

Figure 5.19


12/19


The reason is the very high affinity of chro-

mium to carbon, which causes the precipita-

tion of chromium carbides Cr23C6 on grain

boundaries, Figure 5.22. Due to these precipi-

tations, the austenite grid is depleted of

chromium content along the grain boundaries

and the Cr content drops below the parting

limit. The diffusion speed of chromium in aus-

tenite is considerably lower than that of car-

bon, therefore the chromium reduction cannot

be compensated by late diffusion. In the de-

pleted areas along the grain boundaries (line

2 in Figure 5.22) the steel has become sus-

ceptible to corrosion.

Only after the steel has been subjected to

sufficiently long heat treatment, chromium will

diffuse to the grain boundary and increase the

C concentration along the

grain boundary (line 3 in

Figure 5.22). In this way, the

complete corrosion resis-

tance can be restored (line 4

in Figure 5.22).

Figure 5.23 explains why the

IC is also described as in-

tergranular disintegration.

Due to dissolution of de-

pleted areas along the grain

boundary, complete grains

break-out of the steel.

Carbon Solubility ofAustenitic Cr - Ni Steels

0 0.05 0.1 0.15 0.2 0.25 %0,3Carbon content

600

700

800

900

1000

1100

C

1200

A

Heattreatmenttemperature

to Bain and Aborn


Figure 5.21

Sensibility of a Cr - Steel

Chromium

contentofaustenite

resistance limit

1 - homogenuous starting condition2 - start of carbide formation3 - start of concentration balance4 - regeneration of resistance limit

1

2

3

4

Distance from grain boundary ISF 2002br-er-06-22e.cdr

Figure 5.22


13/19


The precipitation and re-

passivation mechanisms

described in Figure 5.22

are covered by intergranu-

lar corrosion diagrams ac-

cording to Figure 5.24.

Above a certain tempera-

ture carbon remains dis-

solved in the austenite

(see also Figure 5.21).

Below this temperature, a

carbon precipitation takes

place. As it is a diffusion

controlled process, the

precipitation occurs after a

certain incubation time

which depends on tem-

perature (line 1, precipita-

tion characteristic curve).

During stoppage at a con-

stant temperature, the

parting limit of the steel is

regained by diffusion of

chromium.

Figure 5.25 depicts characteristic precipitation curves of a ferritic and of an austenitic steel.

Due to the highly increased diffusion speed of carbon in ferrite, shifts the curve of carbon

precipitation of this steel markedly towards shorter time. Consequently the danger of inter-

granular corrosion is significantly higher with ferritic steel than with austenite.

Grain Disintegration


Figure 5.23

Area of Intergranular Disintegrationof Unstabilized Cr - Steels

Reciprocalofheattreatmenttemperature1/T

oversaturatedaustenite

austenite -chromium carbide (M C )

no intergranular disintegration23 6

unsaturated austenite

Heat treatment time (lgt)

1 incubation time2 regeneration of resistance limit3 saturation limit for chromium carbide

1

2

3

austenite + chromium caride (M C )

to intergranular disintegration23 6 sensitive


Figure 5.24


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As carbon is the element that triggers the intergranular corrosion, the intergranular corrosion

diagram is relevantly influenced by the c con-

tent, Figure 5.26.

By decreasing the carbon content of steel,

the start of carbide precipitation and/or the

start of intergranular corrosion are shifted

towards lower temperatures and longer

times. This fact initiated the development of

so-called ELC-steels (Extra-Low-Carbon)

where the C content is decreased to less

than 0,03%

During welding, the considerable influence of

carbon is also important for the selection of

the shielding gas, Figure 5.27. The higher the

CO2-content of the shielding gas, the

stronger is its carburising effect. The C-

content of the weld metal increases and the

steel becomes more susceptible to inter-

granular corrosion.

An often used method to

avoid intergranular corro-

sion is a stabilisation of the

steel by alloy elements like

niobium and titanium, Fig-

ure 5.28. The affinity of

these elements to carbon is

significantly higher than

that of chromium, therefore

carbon is compounded into

Nb- and Ti-carbides. Now

carbon cannot cause anychromium depletion. The

Precipitation Curves of VariousAlloyed Cr Steels

Tempering time

Temp

eringtemperature

quenchtemperature

18-8-Cr-Ni steel17% Cr steel

precipitation curves for

cooling curve


Figure 5.25

Figure 5.26

Influence of C-Contenton Intergranular Disintegration

101

102

103

104

105

106

Times

400

500

600

700

800

900

1000

Temperatu

re

C

0.07%C0.05%C

0.03%C

0.025%C



15/19


proportion of these alloy elements depend on the carbon content and is at least 5 times

higher with titanium and 10 times higher with niobium than that of carbon. Figure 5.28 shows

the effects of a stabilisation in the intergranular corrosion diagram. If both steels are sub-

jected to the same heat treatment (1050C/W means heating to 1050C and subsequent wa-

ter quenching), then the area of intergranular corrosion will shift due to stabilisation to

significantly longer times. Only with a much higher heat treatment temperature the inter-

granular corrosion accelerates again. The cause is the dissolution of titanium carbides at suf-

ficiently high temperature. This carbide dissolution causes problems when welding stabilised

steels. During welding, a narrow area of the HAZ is heated above 1300C, carbides are dis-

solved. During the subsequent cooling and the high cooling rate, the carbon remains dis-

solved.

If a subsequent stress relief treatment around 600C is carried out, carbide precipitations on

grain boundaries take place again. Due to the large surplus of chromium compared with nio-

bium or titanium, a partial chromium carbide precipitation takes place, causing again inter-

Influence of Shielding Gason Intergranular Disintegration

S hie ld ing g as A r [% ] C O2 O2

S 1 99 / 1

M 1 90 5 5

M 2 82 18 /

Composit ion

0,2 0,5 1 2,5 5 10 25 50 100 250 h 1000400

450

500

550

600

C

700

0.058 % C0.53 % NbNb/C = 9

0.030 % C0.51 % NbNb/C = 17 0.018 % C

0.57 % NbNb/C = 32M2

M1

S1

Heat treatment time



Figure 5.27

Influence of Stabilizationon Intergranular Disintegration

800

700

650

600

550

500

450

C


0,3 1 3 10 30 100 300 1000 h 10000Time

1050C/W

X5CrNi18-10 unstabilized

800

700

650

600

550

500

450

C


0,3 1 3 10 30 100 300 1000 h 10000Time

1300C/W

1050C/W

X5CrNiTi18-10 stabilized

W.-No.:4301 (0,06%)

W.-No.:4541


Figure 5.28


16/19


granular susceptibility. As this susceptibility is limited to very narrow areas along the welded

joint, it was called knife-line attack because of its appearance. Figure 5.29.

In stabilised steels, the chromium carbide represents an unstable phase, and with a suffi-

ciently long heat treatment to transform to NbC, the steel becomes stable again. The stronger

the steel is over-stabilised, the lower is the tendency to knife-line corrosion.

Nowadays the importance

of Nickel-Base-Alloys in-

creases constantly. They

are ideal materials when it

comes to components

which are exposed to spe-

cial conditions: high tem-

perature, corrosive attack,

low temperature, wear re-

sistance, or combinations

hereof. Figure 5.30 shows

one of the possible group-

ing of nickel-base-alloys.

Materials listed there are selected examples, the total number of available materials is many

times higher.

Group A consists of nickel

alloys. These alloys are

characterized by moderate

mechanical strength and

high degree of toughness.

They can be hardened only

by cold working. The alloys

are quite gummy in the an-

nealed or hot-worked con-

dition, and cold-drawn

material is recommended

for best machinability andsmoothest finish.

Knife-Line Corrosion

br-er-06-29e.cdr

Figure 5.29


Alloy Chem. composition Alloy Chem. Composition

Group A Group D1

Nickel 2 00 Ni 99.6, C 0.08 Duranickel 3 01 Ni 94.0, A l 4.4, W 0.6

Nicke l212 Ni 97.0 ,C 0 .05,Mn2.0 Incoloy925 Ni 42.0 ,Fe 32.0 ,Cr 21.0 ,Mo 3 .0 ,W 2.1 ,Cu2.2 ,Al0 .3

Nickel 222 Ni 99.5, Mg 0.075 Ni-Span-C 902 Y2O30.5, Ni 42.5,Fe 49.0,Cr 5.3,W 2.4,Al 0.5

Group B Group D2

Monel 400 Ni 66.5, C u 31.5 Monel K -500 Ni 65.5, C u 29.5, A l 2.7, F e 1.0, W 0.6

Monel450 Ni 30.0,Cu68.0,Fe 0.7,Mn0.7 Inconel718 Ni 52.0,Cr 22.0,Mo 9.0,Co 12.5,Fe 1.5,Al1.2

Ferry Ni 45.0, C u 55.0 Inconel X -750 Ni 61.0, C r 21.5, Mo 9.0, Nb 3.6, F e 2.5

Group C Nimonic 90 Ni 77.5,Cr 20.0,Fe 1.0,W 0.5,Al 0.3,Y2O30.6

Incone l600 Ni 76.0 ,Cr 15.5 ,Fe 8 .0 Nimonic 105 Ni 76.0 ,Cr 19.5 ,Fe 112.4 ,Al1 .4

Ni moni c 75 Ni 80.0, C r 19.5 Incoloy 9 03 Ni 39.0, Fe 34.0, C r 18.0, M o 5.2, W 2.3, A l 0.8

Nimonic 86 Ni 64.0,Cr 25.0, Mo 10.0, Ce 0.03 Incoloy909 Ni 58.0, Cr 19.5,Co 13.5, Mo 4.25, W 3.0,Al 1.4

Incoloy800 Ni 32.5,Fe 46.0, Cr 21.0,C 0.05 Inco G-3 Ni 38.4, Fe 42.0,Cu 13.0,Nb 4.7,W 1.5,Al0.03,Si 0.15

Incoloy825 Ni 42.0,Fe 30.0, Cr 21.5,Mo 3.0, Cu2.2,Ti 1.0 Inco C-276 Ni 38.4,Fe 42.0,Cu13.0,Nb 4.7, W 1.5, Al0.03,Si 0.4

Inco 330 Ni 35.5,Fe 44.0,Cr 18.5,Si 1.1 Group E

MonelR-405 Ni 66.5,Cu 31.5,Fe 1.2, Mn1.1, S 0.04

Typical Classification of Ni-Base Alloys

Figure 5.30


17/19


Group B consists mainly of those nickel-copper alloys that can be hardened only by cold

working. The alloys in this group have higher strength and slightly lower toughness than

those in Group A. Cold-drawn or cold-drawn and stress-relieved material is recommended for

best machinability and smoothest finish.

Group C consists largely of nickel-chromium and nickel-iron-chromium alloys. These alloys

are quite similar to the austenitic stainless steels. They can be hardened only by cold working

and are machined most readily in the cold-drawn or cold-drawn and stress-relieved condition.

Group D consists primary of age-hardening alloys. It is divided into two subgroups:

D 1 Alloys in the non-aged condition.

D 2 Aged Group D-1 alloys plus several other alloys in all conditions.

The alloys in Group D are characterized by high strength and hardness, particularly when

aged. Material which has been solution annealed and quenched or rapidly air cooled is in the

softest condition and does machine easily. Because of softness, the non-aged condition is

necessary for trouble free drilling, tapping and all threading operations. Heavy machining of

the age-hardening alloys is best accomplished when they are in one of the following condi-

tions:

1. Solution annealed

2. Hot worked and quenched or rapidly air cooled

Group E contains only one material: MONEL R-405. It was designed for mass production of

automatically machined screws.

Due to the high number of possible alloys with different properties, only one typical material

of group D2 is discussed here: Material No. 2.4669, also known as e.g. Inconel X-750.

The aluminium and titanium containing 2.4669 is age-hardening through the combination of

these elements with nickel during heat treatment: gamma-primary-phase (') develops which

is the intermetallic compound Ni3(Al, Ti).

During solution heat treatment of X-750 at 1150C, the number of flaws and dislocations in

the crystal is reduced and soluble carbides dissolve. To achieve best results, the material


18/19


should be in intensely worked condition before heat treatment to permit a fast and complete

recrystallisation. After solution heat treatment, the material should not be cold worked, since

this would generate new dislocations and affect negatively the fracture properties.

The creep rupture resistance of X-750 is due to an even distribution of the intercrystalline 'phase. However, fracture properties depend more on the microstructure of the grain bounda-

ries. During an 840C stabilising heat treatment as part of the triple-heat treatment, the fine '

phase develops inside the grains and M23C6precipitates onto the grain boundaries. Adjacent

to the grain boundary, there is a ' depleted zone. During precipitation hardening (700C/20

h) ' phase develops in these depleted zones. ' particles arrest the movement of disloca-

tions, this leads to improved strength and creep resistance properties.

During the M23C6transformation, carbon is stabilised to a high degree without leaving chro-

mium depleted areas along the grain boundaries. This stabilisation improves the resistance

of this alloy against the attack of several corrosive media.

With a reduction of the precipitation temperature from 730 to 620C as required for some

special heat treatments additional ' phase is precipitated in smaller particles. This en-

hances the hardening effect and improves strength characteristics.

Further metallurgical discussions about X-750, can be taken from literature, especially with

view to the influence of heat treatment on fracture properties and corrosion behaviour.

The recommended processes for welding of X-750 are tungsten inert gas, plasma arc, elec-

tron beam, resistance, and pressure oxy arc welding.

During TIG welding of INCONEL X-750, INCONEL 718 is used as welding consumable. Joint

properties are almost 100% of base material at room temperature and about 80% at 700 -

820C. Figure 5.31 shows typical strength properties of a welded plate at a temperature

range between -423 and 1500F (-248 820C).

Before welding, X-750 should be in normalised or solution heat treated condition. However, it

is possible to weld it in a precipitation hardened condition, but after that neither the seam nor

the heat affected zone should be precipitation hardened or used in the temperature range of

precipitation hardening, because the base material may crack. If X-750 was precipitation

hardened and then welded, and if it is likely that the workpiece is used in the temperature

range of precipitation hardening, the weld should be normalised or once again precipitation

hardened. In any case it must be noted that heat stresses are minimised during assembly or

welding.


19/19


X-750 welds should be solution heat treated before a precipitation hardening. Heating-up

speed during welding must be from the start fast and even touching the temperature range of

precipitation hardening only as briefly as possible. The best way for fast heating-up is to in-

sert the welded workpiece into a preheated furnace.

Sometimes a preheating before welding is advantageous if the component to be welded

has a poor accessibility, or the welding is complex, and especially if the assembly proves to

be too complicated for a post heat treatment. Two effective welding preparations are:

1. 1550F/16 h, air cooling

2. 1950F/1 h, furnace cooling with 25-100F/h up to 1200F, air

A repair welding of already fitted parts should be followed by a solution heat treatment (with a

fast heating-up through the temperature range of precipitation hardening) and a repeated

precipitation hardening.

A cleaning of intermediate layers must be carried out to remove the oxide layers which are

formed during welding. (A complete isolation

of the weld metal using gas shielded proc-

esses is hardly possible). If such films are not

removed on a regular basis, they can become

thick enough to cause material separations

together with a reduced strength. Brushing

with wire brushes only polishes the surface,

the layer surface must be sand-blasted or

ground with abrasive material. The frequency

of cleaning depends on the mass of the de-

veloped oxides. Any sand must be removed

before the next layer is welded.

X-750 can be joined also by spot-, projection-,

seam-, and flash butt welding. The welding

equipment must be of adequate performance.

X-750 is generally resistance welded in nor-

malized or solution heat treated condition.

Figure 5.31