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Nickel-Based Superalloy Welding Practices for Industrial Gas Turbine Applications
M.B. Henderson
+
, D. Arrell
b
, M. Heobel*, R. Larsson
b
and G. Marchant
c
+
ALSTOM Power Technology Centre, Whetstone, UK
b
ALSTOM Power Sweden AB, Finspång, Sweden
c
ALSTOM Power (UK) Ltd., Lincoln, UK
*
ALSTOM Power Technology Centre, Daettwil, CH.
Abstract
The continued drive for increased efficiency, performance and reduced costs for industrial gas turbine
engines demands extended use of high strength-high temperature capability materials, such as nickel
based superalloys. To satisfy the requirements of the component design and manufacturing engineers
these materials must be capable of being welded in a satisfactory manner. The present paper
describes the characteristic defects found as a result of welding the more difficult, highly alloyed
materials and reviews a number of welding processes used in the manufacture and repair of nickel
alloy components. These include gas tungsten arc (GTA) and electron beam (EB) welding, laser
powder deposition and friction welding. Many of the more dilute nickel based alloys are readily
weldable using conventional GTA processes, however, high strength, precipitation hardened materials
are prone to heat affected zone and strain age cracking defect formation. A number of factors are
found to affect the propensity for defects: composition (aluminium and titanium content), grain size,
pre and post-weld heat treatment, as well as the welding process itself (control of heat input and
traverse speed). Process parameter identification is still largely empirical and a fuller understanding of
the joining processes is dependent upon the development and application of more sophisticated
numerical modelling techniques.
1. Introduction
Nickel based superalloys are used within the industrial gas turbine (IGT) engine manufacturing
industry, specifically to meet the needs of the hot gas path components. These are exposed to the
most severe operating conditions where high temperature creep, tensile strength, ductility and
oxidation resistance are required to withstand the loadings imposed. A range of nickel based
superalloys, from dilute, solid solution strengthened alloys to the highly alloyed precipitation hardened
materials, have been developed to meet the needs for high temperature structural performance and
environmental resistance.
Figure 1. ALSTOM Power (UK) Cyclone gas turbine engine.
Cost effective and successful manufacture of a modern, high performance IGT engine (see Figure 1)
is dependent upon the ability to join the nickel based superalloy components using methods, such as
gas tungsten arc (GTA), electron beam (EB) and laser welding, and methods such as friction or inertia
bonding. Weldability is defined as “the capacity of a material to be joined under the imposed
fabrication conditions into a specific, suitably designed structure and to perform satisfactorily in the
intended service” [1] providing fitness for purpose with minimal distortion and a controlled/limited
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numbers of defects. These procedures must be capable of being conducted in a cost-effective manner
and the process parameters are largely determined using empirical methods, although more
sophisticated finite element models are becoming available [2] that are capable of simulating the
distortion and formation of a number of characteristic defects. The development of defects within the
weld metal itself and the adjacent heat affected zone (HAZ) of the parent material is dependent upon a
range of factors associated with the joint design and microstructure, form and composition of the
parent alloy and any filler wire used during the welding procedure. The development of excessive
distortion and residual stresses needs to be controlled by suitable assembly jigging, use of optimum
heat input and traverse speed, as well as pre and post-weld heat treatment procedures, if necessary.
Increasingly to achieve through-life cost reduction targets, component refurbishment, overhaul and
repair are a prime consideration for both manufacturers and operators of industrial gas turbines.
Traditionally, refurbishment and repair of a number of dilute nickel alloys has been conducted using
GTA welding procedures and these methods are well established within the industry. However, the
high cost of near-net shape castings, such as turbine blades, vanes and casings, either for first part
manufacture or part-life refurbishment, has necessitated the development and introduction of a
number of more novel joining processes such as laser powder deposition and friction welding.
2. Superalloy Welding and Characteristic Defects
Successful first part manufacturing of a range of components, such as combustor liners (see Figure
2a), transition ducts (see Figure 2b) and exhaust exit casings, is dependent upon the ability to form
fairly complex structures from wrought alloys, such as Nimonic 75, Haynes 230, IN625 and C263.
Likewise, component assembly is dependent upon the ability to perform high quality-high integrity
weldments, generally, for conventional solid solution strengthened alloys using conventional gas
tungsten arc (GTA) welding procedures. These manufacturing methods are well established within the
IGT manufacturing industry and are applied fairly routinely to solid solution and low volume fraction
precipitate strengthened alloys, as mentioned above. Though routine, this technology is key to
achieving the cost and reliability targets specified by the manufacturers and operators, alike. Care
must be taken, however, with coarse grain size materials, which for certain applications have been
heat treated to improve resistance to creep. Successful joining of coarse grained dilute materials is
often restricted to power beam methods such as laser and electron beam welding that introduce lower
thermal transients across adjacent to the weld bead.
Figure 2. a) ALSTOM Power G30 range of lean burn DLE industrial gas turbine combustors.
b) ALSTOM Power industrial gas turbine Combustor transition duct.
Increasingly, there is a requirement to conduct high integrity, defect-free (or limited) welding of higher
temperature capability, precipitation strengthened alloys such as IN718, Waspaloy, IN939, IN738 etc,
as well as more advanced alloys such as those used for directionally solidified and single crystal
castings, such as MarM247 and CMSX-4. These are used for the manufacture of a wide range of hot
gas path components such as discs, casings, stator vane segments and turbine blades. Welding of
these alloys presents much more of a problem due to their highly alloyed nature and the more
complex precipitation strengthening mechanisms needed to provide high temperature strength in
a b
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service. The inherent capabilities of these alloys (ie., the strengthening mechanisms) often interact in
a detrimental manner with the thermal and mechanical loadings generated by the heat source,
component mass and jigging constraints applied during the welding procedure. Local changes in
microstructure adjacent to the weld bead, such as particle coarsening and dissolution and grain
growth within the HAZ can lead to significant property changes during the heating cycle, which interact
with the “thermal fight” arising from precipitation within the alloy during cooling. In an effort to minimise
these interactions many higher strength alloys are welded in the solution annealed or softened
condition. However, despite these efforts the more highly strengthened alloys continue to be
susceptible to three main types of cracking and defects. These are summarised as follows:
• Solidification cracking
Solidification cracking, such as shown in Figure 3a, occurs within the newly formed weld bead when
the mushy, two-phase liquid-solid region experiences tensile stresses and the high fraction of solid
present (typically fs>0.9) restricts the flow of liquid metal to backfill the interdendritic regions. These
are torn apart by tensile thermal stresses generated behind the weld bead as it progresses.
Solidification crack formation is dependent upon a number of contributory factors such as thermally
induced stresses and strains being generated behind the weld bead, that coincide with a high fraction
of solid being present in the mushy zone, solidification and microsegregation, viscous flow of liquid
metal and crack initiation and propagation effects. Formation is promoted by a wide solidification
range for the alloy (i.e., dilute and eutectic forming alloys are less susceptible) and low welding
traverse speeds that promote the generation of tensile stresses adjacent to the weld due to
contraction of the surrounding solid material. This form of welding defect in nickel based alloys can,
generally, be avoided by optimising the welding procedures used. A fuller discussion of the factors
leading to solidification cracking and the analytical and finite element modelling methods used to
simulate these phenomena has been given elsewhere [3].
Another common type of defect found during welding of thin plate superalloy components, that is often
associated with solidification cracking, is a continuous grain boundary that forms along the centreline
of the weld bead at intermediate to high heat input levels and high traverse speeds (see Figure 3b). A
fuller description of the formation of centreline grain boundary and the dendrite-tip growth kinetics
model capable of predicting its formation have been given elsewhere [4]. Formation of this defect is
promoted by higher alloying additions and impurity levels and is characterised by a sharp teardrop-
shaped weld pool, a coarse columnar grain structure across the weld bead and a high volume fraction
of eutectic and brittle phases along the centreline. This segregation may lead to incipient melting
during subsequent heat treatment and localised corrosion of the weld bead during service exposure.
Figure 3. a) Solidification cracking following autogenous GTA welding of IN718 sheet. [5]
b) Centreline grain boundary formation in IN718 sheet.
• Grain boundary liquation cracking
Grain boundary liquation cracking or heat affected zone (HAZ) fissuring, as shown in Figure 4 for GTA
welding of IN718, occurs within the HAZ adjacent to the weld bead as a consequence of local
dissolution of grain boundary phases, such as primary MC and M
6
C carbides, Laves phases and σ-
phase [6, 7, 8, 9]. Under rapid heating, the grain boundary phases are unable to dissolve fully into the
surrounding matrix and partial dissolution leads to the formation of a low melting point eutectic and
melting of the grain boundary region. A liquid film forms on grain boundaries within the HAZ, often
away from the fusion zone in regions where grain coarsening has occurred, which fails under the
tensile thermal stresses generated immediately behind the weld bead.
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Grain boundary liquation has been found to be the primary cause of HAZ hot cracking in alloys such
as IN718 [9] and is associated with grain boundaries rich in those elements that form primary MC-type
carbides (NbC and TiC in IN718, for example). Susceptibility to HAZ fissuring is dependent upon alloy
composition (carbon, boron etc), grain size and grain boundary character [6], and for coarse grain size
materials (> ASTM 6), the weld traverse speed. Increasing traverse speed increases the HAZ cracking
susceptibility as this influences the thermal gradients and stress state within the HAZ. Generally, fine
grained wrought materials, such as wrought IN718, are less susceptible and are considered to be
readily weldable provided suitable pre and post-weld heat treatments are applied; typically the material
is welded in the solution annealed condition. An increased grain size decreases the grain boundary
area per unit volume and increases the amount of segregation on the grain boundaries. In addition, it
is clear that the segregation of tramp and impurity elements such as boron, sulphur and phosphorus
also plays a role [10, 11] and is promoted by coarser microstructures. These elements act to suppress
the eutectic melting temperature and consequently increase the HAZ fissuring susceptibility, however,
it appears to be possible to control their segregation behaviour with suitable solution heat treatment
procedures [6, 7].
Figure 4. Grain coarsening and HAZ liquation cracking at A following autogenous GTA welding of
IN718 sheet. [5]
Attempts have been made to improve the HAZ cracking resistance of high temperature casting alloys,
such as IN939, by reducing the impurity levels present in the alloys and refining the final
microstructure [12]. However, the coarse grain size microstructures that are typically found for
investment cast or near net shaped cast parts are generally found to be susceptible to HAZ cracking.
It is evident that for a number of fine grained, precipitation strengthened alloys, such as IN718 and
Waspaloy, it is possible to produce high quality welds without HAZ cracking problems by controlling
the alloy composition, grain size, weld traverse speed and heat input by using EB welding methods.
However for a wide range of materials, in particular the high strength cast alloys such as IN738,
IN939, MarM247 and cast forms of IN718, it has, generally, not been possible within a production
environment to avoid HAZ fissuring using conventional GTA and EB welding methods.
• Strain age cracking
Strain age or re-heat cracking, generally, occurs in γ’-Ni
3
(Al, Ti) precipitate strengthened alloys during
post-weld heat treatment or subsequent high temperature service due to the presence of either
residual stresses developed during manufacture, or applied stresses arising from service exposure.
These defects are characterised by intergranular micro-cracking in either the HAZ or weld bead and
form as a consequence of precipitation and hardening of the alloy during thermal exposure and
transfer of solidification strains onto the grain boundaries [10, 13], often with carbides acting as crack
initiation sites. It is common practice to attempt to relieve residual stresses arising from the welding
procedure by means of post weld heat treatment. However, often the stress relieving temperature is
greater than the ageing temperature of the alloy and this leads to a transient precipitation period
during post weld heating that hardens the alloy and leads to excessive strain localisation on grain
boundaries within the HAZ and weld bead during heating. Conventionally, the most effective means of
limiting the extent of strain age cracking is to overage the material prior to welding. This can be
combined with the use a more ductile, dilute alloy filler wire (for example, IN625 or C263) and by
careful control of the heating and cooling cycles during post-weld heat treatment [10]. However,
welding of more difficult alloys in the fully solutioned or over aged condition can lead to HAZ cracking
due to re-precipitation of strengthening phases during cooling immediately after welding and thus lead
to cracking sensitisation.
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3. Weldability Assessment Diagrams
The susceptibility to strain age cracking is promoted by high additions of both Ti and Al, as these
promote γ’ precipitation, and elements such as carbon, sulphur and boron. Alloys such as IN718 are
considered as being less susceptible due to the more sluggish precipitation reaction of γ’’-Ni
3
(Nb, Ti,
Al, Mo) precipitates. The weldability of nickel based superalloys and the susceptibility to strain age
cracking is often assessed, qualitatively, by plotting the Al vs. Ti content of the alloy, as shown in
Figure 5. When the total Al+Ti level for a particular alloy exceeds a critical value (often taken as 4wt%)
it is deemed to be difficult to weld and increasingly unweldable with increasing Al+Ti content. This over
simplification of the factors contributing towards weld cracking does not take any account of variation
in microstructure due to different thermomechanical processing and heat treatment procedures. Those
alloys that lie either side of the critical 4wt% limit need to be treated carefully in terms of the heat
treatment, grain size and cooling rates used.
Weldability Assessment: strain age cracking susceptibility
C263
CMSX-4
IN738LC
B1900IN713C
R'108
Mar-M-200
IN100
AF2-1DA
Astroloy
Udimet 700
Udimet 600
Udimet 500
GMR 235
Inconel 700
IN939
Rene'41
Unitemp 1753
Waspaloy
Rene'62
M252
Inconel X-750
Inconel X
Inconel 909
Rene'220C
Inconel 718 (x)
0
1
2
3
4
5
6
7
0 1 2 3 4 5 6 7
Ti content, wt-%
Al content, wt-%
Difficult to weld: weld and
strain age cracking
Weldable
Figure 5. Weldability assessment diagram for a range of nickel based superalloys (after[10]).
Recent weld research programmes have been aimed at establishing a more systematic approach to
evaluating and predicting the weldability of nickel based superalloys in terms of the process
parameters (i.e., heat/power input and traverse speed) applied during the welding procedure for a
particular alloy, welding process and joint configuration. Figure 6 shows a typical Weldability Process
Diagram for wrought sheet IN718, which illustrates the defect formation regimes as discussed above,
which have been identified by a combination of analytical and numerical modelling techniques in
conjunction with well characterised experimental validation of the weld defect formation zones. The
limits of these boundaries identify a central region of weldability for this alloy. With the advent of more
sophisticated, non-linear finite element modelling and data analysis tools it should be possible to
identify similar weldability diagrams for a range of alloys that are currently considered to be difficult to
weld and provide a more comprehensive tool to define weldability limits for nickel based superalloys.
4. Nickel Based Superalloy Component Welding Practices
4.1 Conventional Welding and Repair Procedures Using GTA
As discussed previously, conventional, manual and automatic GTA welding procedures are used in
the production of complex shaped components, such as the transition duct shown in Figure 7. This
unit is manufactured in the solid solution and carbide precipitate strengthened alloy Haynes 230 sheet
and forged plate. GTA methods are applied routinely, with little difficulty concerning the integrity of the
joint (see Figure 8a). The main issues for this component are associated with the high cost of
![Figure 6. GTA Weldabilty Process Diagram for IN718 plate (after [5]).](/figures/figure-6-gta-weldabilty-process-diagram-for-in718-plate-30s3h323.webp)
