SupeR ALLOYS PART 2
Properties of Nickel based Superalloys
Superalloys
constitute a large fraction of the materials of construction in turbine engines
because of their unique combination of physical and mechanical properties.
Table 7 lists some typical physical properties of superalloys. In aircraft
engines, it is typical to consider density-normalized properties; thus alloy
densities, which are typically in the range of 7.7–9.0 g/cm3, are of specific
interest. Optimization of the relevant set of mechanical properties is of
paramount importance and is dependent on a high level of control and
understanding of the processes, because mechanical properties are a strong
function of microstructure. Mechanical properties of primary interest include
tensile properties. creep, fatigue, and cyclic crack growth.
Depending on the
details of component design, any one of these four properties may be life
limiting. Each of these properties are briefly discussed in the following
sections. It is worth noting that models for prediction of these properties
must treat many aspects of microstructure at different length scales.
Table 7.Typical
physical properties of superalloys
|
Property
|
Typical ranges
|
|
Density
|
7.7–9.0 g/cm3
|
|
Melting temperature (liquidus)
|
1320–1450◦C
|
|
Elastic modulus
|
Room temp: 210 GPa
800◦C: 160 GPa
|
|
Thermal expansion
|
8–18×10−6/◦C
|
|
Thermal conductivity
|
Room temp: 11 W/m ·K
800◦C: 22 W/m·K
|
Tensile Properties
Nickel-based
superalloys have relatively high yield and ultimate tensile strengths, with
yield strengths often in the range of 900–1300 MPa and ultimate tensile
strengths of 1200–1600MPa at room temperature. Turbine-disk alloys are
typically developed to have higher strengths for flexibility at temperatures
below 800◦C in the design to protect against burst of the turbine disk in the
event of an engine over speed. Note that the tensile properties do not
substantially decay until temperatures are greater than approximately 850◦C.
The slight rise in the yield strength of the alloys at intermediate
temperatures is due to the unusual flow behavior of the Ni3Al γ’ phase. The
temperature dependence of yielding in single-phase Ni3Al, with a strong
increase at intermediate temperatures, is shown in Fig. 7.
Deformation
of the precipitates gives a corresponding, but weaker, rise in the flow stress
of superalloys at intermediate temperatures. Note also that the two-phase
superalloys are much stronger than either the matrix or precipitate materials
in their bulk form figure given below. Strengthening in two-phase superalloys
arises from multiple microstructural sources, including solid-solution
strengthening, grain size strengthening, and the interaction of dislocations
with precipitates (Orowan bowing between precipitates or shearing through
precipitates in strong- or weak-coupled modes).
It
is worth noting that heat treatments and
various processing steps at different temperatures can result in
multiple populations of precipitates with substantially varying mean sizes.12For
example, three distinct populations of precipitates in powder IN100 are
apparent in Fig. 17, where the microstructure is imaged at various
magnifications. The three populations of
precipitates are referred to as
primary, secondary, and tertiary, in order of decreasing size and
precipitation temperature. The primary precipitates are not solutioned in the
late stages of processing and therefore inhibit grain growth and indirectly
provide grain size strengthening.
Fig.7 Comparison of
the critical resolved shear stress (CRSS) corresponding to the Ni-based
superalloy MAR-M200 and the individual constituent phases.
Fig. 8 Microstructure
of subsolvus and supersolvus processed IN100
B. Creep Properties
Because
superalloys experience extended periods under stress at high temperature, a
high resistance to time-dependent creep deformation is essential. This is very
important for cast blade alloys, because they will experience temperatures up
to 1100°C, whereas disk alloys are typically limited to less than 700°C. For a
fixed stress and temperature, two-phase superalloys have a much higher creep
resistance compared to their single-phase counterparts (Fig. 9). As with all
properties that are governed by plastic deformation processes, creep properties
are sensitive to microstructure.
Fig.9 As reflected by the minimum creep rates,
two-phase γ–γ_ superalloys exhibit significantly improved creep resistance when compared
to their single-phase constituents.
Figure
10 shows the creep-rupture life as a function of volume fraction of
precipitates for a single crystal alloy.57 Note that the strength peaks when
the precipitate volume fractions are in the range of 0.6–0.7. Not surprisingly,
many alloys contain volume fractions of precipitates in this peak range. Alloy
chemistry is also important to creep properties. Because the rate-controlling
processes in creep are diffusion controlled, elements that have low
interdiffusion coefficients with nickel are generally beneficial to creep.
Interdiffusion for various alloying elements in nickel has recently been
studied in detail.58,59 Elements most effective at slowing diffusion include
Ir, Re, Ru, Pt, W, Rh, and Mo.
Fig.10 Variation
in creep rupture life as a function of γ_ volume fraction for single crystal TMS-75
A
combination of increasing refractory alloying additions and advances in
processing has resulted in substantial increases in the maximum temperature
capability of superalloys over the past few decades. For example, considering a
creep-rupture life of 1000 h at a stress of 137 MPa, the most recently
developed single-crystal superalloys have a temperature capability of
approximately 1100°C, whereas conventionally cast equiaxed alloys developed in
the 1970s had a temperature capability of 900–950°C. The temperature
capabilities have reached 85–90% of the melting point, and usually high
fraction of melting, compared to the operating conditions of any other class of
structural materials. This indicates a need for development of a completely new
class of materials with higher melting points; unfortunately this is a major
challenge and there are no obvious replacements for superalloys in the hottest
sections of the turbine engine. The exceptional creep properties are largely
due to the fact that the high resistance of the precipitates to shearing
extends to elevated temperatures. The uniaxial stress σOR required
to glide a dislocation through the narrow matrix channels of the superalloy
microstructure is,
σOR =
(μb/hS)
where
μ is the shear modulus, b is the Burgers vector, h is the width of the channel,
and S is the Schmid factor. Typical values of these material properties at 850◦C
are μ=48.2 GPa, b=0.254 nm, and h =60 nm. For these parameters, an applied
stress of 408 MPa must be exceeded for the onset of dislocation glide through
the channels at 850◦C. Thus, this resistance accounts for a large fraction,
although not all, of the creep resistance of the two-phase material. Figure 11 shows
dislocations gliding through matrix channels during creep of CMSX-3 at 850◦C
and 552MPa. There have been recent efforts to model the creep behavior of
superalloys with the use of continuum damage-mechanics approaches; however,
formulating models that capture the essence of the wide array of complex
deformation mechanisms remains a challenge. The details of the deformation
processes are very sensitive to temperature and applied stress, and it is most
convenient to consider mechanisms of creep deformation at low, intermediate,
and high temperatures (and high, intermediate, and low stresses).
Fig. 11 After an
initial incubation period where dislocations
fill the horizontal channels, dislocations are
forced to bow through the vertical channels
because of the high resistance of the
precipitates to shearing.
1. Low-Temperature Creep
Superalloys
are very resistant to creep deformation at temperatures below 800°C. In general
creep deformation occurs by deformation on <110> {111} slip systems, with
an initial preference for dislocation glide through the continuous matrix.
However, at temperatures <0.6TM, high uniaxial stresses may result in the
activation of <112> {111} slip systems. As dislocations accumulate at the
γ –γ’ interface, ½<110> dislocations undergo reactions that result in
<112> type dislocations that are able to penetrate into the γ’ precipitate.
The details of the dislocation reactions that result in such
precipitate-shearing processes remain under investigation. Nevertheless, as
this shearing process occurs, strain is accumulated rapidly as a result of the
planar nature of the slip and can result in high strains during the primary
creep transient. This mode of deformation is observed in both single crystal
and polycrystalline alloys during creep. For single crystals, [110]- and
[111]-oriented crystals are less prone to this type of deformation, whereas in
[001] crystals, planar slip along <112> continues until crystal rotations
enable the resolved shear stresses to activate slip on other planes. Other
factors that influence these shearing processes include alloy composition, γ’
size, and volume fraction.
2. Intermediate-Temperature Creep
At
intermediate temperatures, stress levels are typically insufficient to result
in shearing of the γ’ precipitates.
Thus, deformation within the microstructure is generally confined to the γ
matrix and results in unusual creep curves that contain an initial incubation
period and a brief primary transient, followed by an extended period of
accelerating creep. In general a steady-state creep rate is not achieved.
No
macroscopic straining occurs during the incubation period, in part due to the
low density of grown-in dislocations. These initial dislocations serve as
sources from which dislocations in the γ matrix are able to multiply.
Single-crystal experiments show that when an external uniaxial stress is
applied, the misfit stresses between the γ matrix and γ’ precipitate are unbalanced and the effective
stresses enable preferential flowof dislocations within the horizontal
channels. The incubation period ends when dislocation percolation, initially through
the horizontal matrix channels and later through the vertical matrix channels,
is complete.
The
deformation mechanism associated with the primary creep transient during
intermediate-temperature creep is distinctly different from that of
low-temperature creep. Unlike dislocation shearing of the γ’ precipitates along the <112> direction at high stresses and low
temperatures, the primary creep transient at intermediate temperatures can be
attributed to the relief of coherency stresses at the γ –γ’ interface as
dislocations accumulate. At the end of primary creep, a three-dimensional
network of dislocations is formed around the precipitates. Despite the lack of
a steady-state strain rate, these dislocation networks surrounding the
precipitates are extremely stable and contribute to the gradual progression
into tertiary creep.
3. High-Temperature Creep
The
enhanced diffusivity associated with deformation at extremely high temperatures
results in morphological changes within the microstructure. With the
application of an external stress to alloys with significant precipitate-matrix
misfit, the discrete cuboidal γ’
precipitates coalescence into rafts or rods aligned perpendicular or
parallel to the applied-stress direction.
The kinetics of
this directional coarsening process are strongly influenced by the temperature
as well as the stresses associated with the coherency strains at the γ –γ
‘
interface. Alignment of the rafts or rods,
however, is dependent upon whether the external and misfit stresses are
compressive or tensile. For example, uniaxial tensile stresses cause alloys
with negative misfit to form rafts perpendicular to the applied-stress
direction, whereas compressive stresses applied to the same alloy will result
in the formation of rods aligned parallel to the direction of applied stress.
Because most commercial directionally solidified and single-crystal alloys
exhibit a negative misfit and are used to sustain tensile loads, rafts are
generally formed perpendicular to the applied-stress direction (Fig. 12).
Fig. 12
Directional coarsening at elevated
temperatures results in the formation of γ rafts aligned perpendicular
to the direction of the applied stress. The matrix inverts as the γ rafts
are contained within a γ’ matrix.
With
single-crystal and directionally solidified Ni-based superalloys containing in
excess of 60% γ’ by volume, directional
coarsening may also result in an inversion of the microstructure. Once rafting
is complete, rafts of γ are contained within an intermetallic γ’ matrix. Because the rafts of γ are discrete,
a continuous path for dislocation motion along the matrix channels no longer
exists. For deformation to continue, dislocations must shear the γ’ phase.
Provided that the rafted structure remains stable, rafted structures are very
resistant to deformation at low stresses where dislocations are unable to
penetrate into the z’ matrix. When
stresses are sufficient to cause shearing of the γ’ phase, microstructural damage is able to accumulate
rapidly and tertiary creep rates are accelerated.
C. Fatigue and Fatigue Crack Growth
Turbine-engine
components experience significant fluctuations in stress and temperature during
their repeated takeoff–cruise–landing cycles. These cycles can result in
localized, small, plastic strains. Thus low-cycle, low-frequency fatigue is of
interest to engine design. Engine vibrations and airflow between the stages of
the turbine can also result in high-cycle fatigue with rapid cycle accumulation
in airfoils at much higher frequencies, in the kHz range.
Figure
13 shows fatigue properties for the powder-disk alloy Rene’ 88 DT at room temperature and 593°C . Such
polycrystalline-disk alloys exhibit outstanding fatigue properties, with
fatigue strengths that are a high fraction of the monotonic yield strength.
Under strain-controlled testing conditions in the low-cycle regime, extensive
cyclic hardening typically occurs. At lower temperatures, cyclic deformation
tends to occur fairly inhomogeneously by shearing of precipitates along {111}
planes, with dislocations isolated in planar slip bands. As the ease of cross
slip increases with temperature, deformation occurs more homogeneously through
the microstructure.
Fig. 13 S-N behavior of Rene’ 88 DT at 20 and 593◦C at a load ratio of 0.05 at frequencies of 10 Hz and 20 kHz. An arrow
indicates a specimen that ran out to 109 cycles.
Fatigue
properties are sensitive to mean stress, particularly at elevated temperatures.
Variability in fatigue lives tends to increase at lower stresses and longer lives
in superalloys. An example of this is shown in Fig. 12 for Rene’ 88 DT, where
fatigue lives at 593°C vary between approximately 106 cycles and 109 cycles at
R =0.05 and σmax =600 MPa. Such variability may be due to intrinsic variations
in microstructure and/or due to the infrequent appearance of extrinsic defects.
Crack initiation during fatigue of disk alloys may occur at extrinsic
inclusions that are introduced during processing or at specific microstructural
features such as larger grains. In cast alloys, cracks may also initiate at
porosity (Fig. 14), carbides, or eutectic. There is a general tendency for
initiation to shift from fatigue sample surfaces to subsurface regions in the
high-cycle regime, for both equiaxed and single-crystal alloys.
Fig.14 Fatigue-crack initiation at a near-surface
pore in a singlecrystal superalloy (photo courtesy of L. Rowland).
Because
of stringent safety and associated lifing requirements, fatigue-crack
propagation is also an important aspect of material behavior, particularly for
turbine-disk materials. Paris law crack growth behavior is displayed over a
wide range of stress intensities for most superalloys. Crack-growth rates are
sensitive to microstructural features, including grain size, precipitate sizes,
and volume fractions. At temperatures above approximately 500°C, environmental
effects and cyclic frequency become significant factors, with higher
crack-growth rates in air than vacuum.
In single-crystal alloys, cracks may grow
crystallographically along {111} planes, particularly in the early stages of
growth (stage I). Depending on testing conditions, as cracks grow longer (stage
II) they may advance in a less crystallographic manner with a greater tendency
toward mode 1 behavior.
Although fatigue
and fatigue-crack growth are often limiting properties, comprehensive models
for life prediction that account for complex loading, crack initiation, and
crack growth as well as microstructure and environment continue to be
developed. Integration of physics-based models with advanced sensors that can
diagnose the current “damage state” remains a promising approach for component
life prediction.
Hot
corrosion resistance
Hot
corrosion resistance was evaluated by crucible test , keeping a piece of alloy
(6-8 mm in diameter and 3-5 mm in height) in a salt mixture (Na2SO4
-25%NaCl) open to air at 900 deg.C for 20 hours. The resistance was
quantitatively specified by metal loss after all the scales being removed.
Morphologically the hot corrosion was classified into three types ; Type I :
corrosion layer composed of Cr sulfide, Ni sulfide, and porous oxide, TypeII :
corrosion layer of thin tight Cr2O3 with a slight or no
amount of sulphide in matrix, Type III : corrosion layer composed of three
layers of oxides, Cr2O3, TiO2, and Al2O3
from outside to inside with a little amount of Cr rich sulphide dispersed in
matrix. A regression analysis was carried out over 42
alloys giving type I corrosion. The resultant equation is shown in Table 9 (For
comparison sake another regression equation is shown in the same table for a
burner rig test conducted elsewhere ; both equations show a good agreement
except for Nb). The equation indicates that Hf doped and a high Cr and Ti
containing alloy is the most preferable in gamma-prime precipitation hardening
alloys, while the addition of W, Ta, or Mo, which are essential for increase of
high temperature strength, is extremely harmful for the hot corrosion. The
equation does not give a good prediction of other types of hot corrosion than
type I. This means that there is(are) other factors) to govern the hot
corrosion. Although it is not clear what they are, our experiment has suggested
that content of Ti plays an important role.
Common nickel base superalloys and
their applications
Inconels
Inconel is a registered trademark of Special Metals Corporation that refers to a family of austenitic nickel-chromium-based superalloys. Inconel alloys are typically used in
high temperature applications. Common trade names for Inconel include: Inconel
625, Chronin 625, Altemp 625, Haynes 625, Nickelvac 625 and Nicrofer 6020.
Typical properties
Inconel alloys are oxidation and corrosion resistant materials well suited for
service in extreme environments subjected to pressure and heat. When heated,
Inconel forms a thick, stable, passivating oxide layer protecting the surface
from further attack. Inconel retains strength over a wide temperature range,
attractive for high temperature applications where aluminum and steel would succumb to creep as a result of thermally-induced
crystal vacancies . Inconel's high temperature strength is developed by solid solution strengthening or precipitation strengthening, depending on the alloy. In age hardening or precipitation strengthening
varieties, small amounts of aluminium combine with nickel to form the intermetallic compound Ni3Al or gamma
prime (γ'). Gamma prime forms small cubic crystals that inhibit slip and creep effectively at elevated
temperatures.
Applications
Inconel is often encountered in
extreme environments. It is common in gas turbine blades, seals, and combustors, as
well as turbocharger
rotors and seals, electric submersible well pump motor shafts, high temperature
fasteners, chemical processing and pressure vessels, heat exchanger tubing, steam generators in nuclear pressurized water reactors, natural gas processing with contaminants such as H2S and CO2, firearm sound suppressor blast baffles, and Formula One and NASCAR
exhaust systems. It is also used in the exhaust systems of high powered rotary
engined Norton motorcycles where exhaust temperatures reach more than 1,000
degrees C. Inconel is increasingly used in the boilers of waste incinerators. Inconel is also used for the
lightweight sport exhaust of recent supercars, the McLaren MP4-12C, the Pagani Huayra and the Koenigsegg Agera.
Inconel is also used in the construction of
higher end firearms sound suppressors and muzzle devices. This is
especially common in suppressors designed to be especially small or for use
with machine guns. Alternatives to the use of Inconel in chemical applications
such as scrubbers, columns, reactors, and pipes are Hastelloy, perfluoroalkoxy
(PFA) lined carbon steel or fiber reinforced plastic.
Hastelloys
Hastelloy is the registered trademark
name of Haynes International, Inc. The trademark is
applied as the prefix name of a range of twenty two different highly corrosion-resistant
metal
alloys loosely grouped by the metallurgical
industry under the material term “superalloys”
or “high-performance alloys”.
The predominant alloying ingredient is typically
the transition metal nickel. Other
alloying ingredients are added to nickel in each of the subcategories of this
trademark designation and include varying percentages of the elements molybdenum,
chromium, cobalt, iron, copper, manganese, titanium, zirconium, aluminum, carbon, and tungsten.
Applications
The primary function of the Hastelloy super
alloys is that of effective survival under high-temperature, high-stress
service in a moderately to severely corrosive, and/or erosion-prone environment
where more common and less expensive iron-based alloys would fail, including
the pressure vessels of some nuclear
reactors, chemical reactors, distillation equipment, and
pipes and valves in chemical industry. Although a super alloy, Hastelloy does
experience degradation due to fabricating and handling. Electropolishing
or passivation of Hastelloy can improve
corrosion resistance.
Monels
Compared to steel, Monel is very
difficult to machine as it work-hardens very quickly. It needs to be turned and
worked at slow speeds and low feed rates. It is resistant to corrosion and
acids, and some alloys can withstand a fire in pure oxygen. It is commonly used
in applications with highly corrosive conditions. Small additions of aluminium and titanium form an alloy (K-500) with the same
corrosion resistance but with much greater strength due to gamma prime
formation on aging. Monel is typically much more expensive than stainless steel.
Marine applications
Monel's corrosion resistance makes it ideal for marine
applications such as piping systems, pump shafts, seawater valves, trolling
wire, and strainer baskets. Housings for magnetic-field measurement equipment.
In recreational boating, Monel wire is used to seize shackles for anchor rodes,
Monel is used for water and fuel tanks, and for under water applications. It is
also used for propeller shafts and for keel bolts.However, because of the problem of electrolytic action in salt water (also known as Galvanic corrosion), in shipbuilding monel must be carefully insulated from other metals such as steel.
Applications in nuclear engineering
The good resistance against corrosion
by acids and oxygen makes monel a good material for the chemical industry. Even
corrosive fluorides can be handled within monel apparatus; this was done in an
extensive way in the enrichment of uranium in the Oak Ridge Gaseous Diffusion Plant. Here most of the larger diameter tubing for the uranium hexafluoride was made from monel. Regulators for reactive cylinder gases like hydrogen chloride form another example, where PTFE is not a suitable option when high delivery pressures are
required. These will sometimes include a Monel manifold and taps prior to the
regulator that allow the regulator to be flushed with a dry, inert gas after
use to further protect the equipment.
RENE 41
Rene 41 is a
precipitation hardening, nickel-based high temperature alloy possessing high
strength in the 1200/1800°F (649/982°C) temperature range. This alloy is
designed for use in severely stressed high temperature applications. Rene 41
has found applications in jet engine and high speed airframe components such
as: afterburner parts, turbine casings, wheels, buckets, bolts, and fasteners.
Rene 41 is highly corrosion and oxidation resistant.
It provides very good resistance to jet engine combustion gases up to 1800°F
(982°C)
Iron- Nickel based superalloys
The iron–nickel based superalloys originally evolved
from austenitic stainless steels with at least 25% nickel required to
stabilize the FCC austenitic matrix.
Most of them contain 25–45% nickel and 15–60% iron. Chromium from15 to 28% is
added for oxidation resistance at elevated temperature, while 1–6% molybdenum
is added for solid solution strengthening. Titanium, aluminum and niobium are added for precipitation hardening. In
the iron–nickel based alloys, the
most common precipitate is γ’ as in
the alloy A-286, which contains 26% nickel.
In alloys containing niobium, such as
Inconel 718, the γ’ Precipitate Ni3Nb is the predominate
strengthener. Due to the lower temperature stability of the
γ’’ precipitate compared to the
ɣ’ precipitate, the maximum use temperature of Inconel 718 is about 1200°F;
however, Inconel 718 is still the most widely used of all of the superalloys.
It is one of the strongest at low temperatures but rapidly loses strength in
the
1200–1500°F range.
The iron–nickel based family also contains some low expansion
alloys, such as Incoloy 903, which are important in applications requiring
closely controlled clearances between rotating and static components. Some iron–nickel alloys are strengthened primarily by
solidsolution hardening, such as
19-9DL, which is essentially 18–8 stainless steelwith slight chromium and nickel adjustments. As a class, the iron–nickel
basedalloys have useful strengths to about 1200°F.
Iron-Based
superalloys
The
iron-based grades, which are less expensive than cobalt or nickel-based grades,
are of three types: alloys that can be strengthened by a martensitic type of
transformation, alloys that are austenitic and are strengthened by a sequence
of hot and cold working (usually, forging at 2,000 to 2,100°F followed by
finishing at 1,200 to 1,600°F), and austenitic alloys strengthened by precipitation
hardening. Some metallurgists consider the last group only as superalloys, the
others being categorized as high-temperature, high-strength alloys. In general,
the martensitic types are used at temperatures below 1,000°F; the austenitic
types, above 1,000°F.
The
AISI 600 series of superalloys consists of six subclasses of iron-based alloys:
·
601 through 604: Martensitic low-alloy steels.
·
610 through 613: Martensitic secondary hardening
steels.
·
614 through 619: Martensitic chromium steels.
·
630 through 635: Semiaustenitic and martensitic
precipitation-hardening stainless steels.
·
650 through 653: Austenitic steels strengthened
by hot/cold work.
·
660 through 665: Austenitic superalloys; all
grades except alloy 661 are strengthened by second-phase precipitation.
·
Iron-based
superalloys are characterized by high temperature as well as room-temperature
strength and resistance to creep, oxidation, corrosion, and wear. Wear
resistance increases with carbon content. Maximum wear resistance is obtained
in alloys 611, 612, and 613, which are used in high-temperature aircraft
bearings and machinery parts subjected to sliding contact. Oxidation resistance
increases with chromium content. The martensitic chromium steels, particularly
alloy 616, are used for steam-turbine blades.
The
superalloys are available in all conventional mill forms -- billet, bar, sheet,
and forgings -- and special shapes are available for most alloys. In general,
austenitic alloys are more difficult to machine than martensitic types, which
machine best in the annealed condition. Austenitic alloys are usually
"gummy" in the solution-treated condition and machine best after
being partially aged or fully hardened. Crack sensitivity makes most of the
martensitic steels difficult to weld by conventional methods. These alloys
should be annealed or tempered prior to welding; even then, preheating and
postheating are recommended. Welding drastically lowers the mechanical
properties of alloys that depend on hot/cold work for strength.
All
of the martensitic low-alloy steels machine satisfactorily and are readily
fabricated by hot working and cold working. The martensitic secondary-hardening
and chromium alloys are all hot worked by preheating and hot forging.
Austenitic alloys are more difficult to forge than the martensitic grades.
Properties
of Iron based Superalloy A-286
Type A-286 alloy (S66286) is an iron-base superalloy useful for
applications requiring high strength and corrosion resistance up to 1300°F
(704°C) and for lower stress applications at higher temperatures.
Type A-286 alloy is a heat and corrosion
resistant austenitic iron base material which can be age hardened to a high
strength level. The alloy is also used for low temperature applications
requiring a ductile, non-magnetic high strength material at temperatures
ranging from above room temperature down to at least -320°F (-196°C). The alloy
may be used for moderate corrosion applications in aqueous solutions.
Type A-286 alloy can be produced by AOD
refining or vacuum induction melting. Vacuum arc or electroslag remelting
procedures may be used to further refine the material. Type A-286 alloy is available in plate, sheet and strip.
Chemical
Composition of Superalloy A-286
Typical chemical analysis of Super Alloy A-286
|
Element
|
Percentage
|
|
Carbon
|
0.08 max
|
|
Manganese
|
0.35 max
|
|
Phosphorus
|
0.015 max
|
|
Sulfur
|
0.015 max
|
|
Silicon
|
0.35 max
|
|
Chromium
|
17-21
|
|
Nickel
|
50-55
|
|
Molybdenum
|
2.80-3.30
|
|
Columbium
|
4.75-5.50
|
|
Titanium
|
0.65-1.15
|
|
Aluminum
|
0.20-0.80
|
|
Cobalt
|
1.00 max
|
|
Boron
|
0.006 max
|
|
Copper
|
0.30 max
|
|
Tantalum
|
0.05 max
|
|
Iron
|
Balance
|
Corrosion
and Oxidation Resistance of A-286
Type A-286 alloy content is similar in
chromium, nickel, and molybdenum to some of the austenitic stainless steels.
Consequently, A-286 alloy possesses a level of aqueous
corrosion resistance comparable to that of the austenitic stainless steels. In
elevated temperature service, the level of corrosion resistance to atmospheres
such as those encountered in jet engine applications is excellent to at least
1300°F (704°C).
Oxidation resistance of A-286 is high for continuous service up to 1500°F (816°C) and
intermittent service up to 1800°F (982°C).
Physical Properties
of A-286
Typical physical properties of A-286 are
summarized in the table below.
|
|
Solution Treated
|
Solution Treated and Aged
|
|
Density
|
0.286 lb/in3 (7.92 g/cm3)
|
0.287 lb/in3 (7.94 g/cm3)
|
|
Specific Gravity
|
7.92
|
7.94
|
|
Melting Range
|
2500-2600°F
|
1370-1430°C
|
Linear
Co-Efficient of Thermal Expansion of A-286
The linear co-efficient of thermal expansion for solution treated
and aged A-286 is provided below.
|
Temperature Range
|
Linear Co-Efficient of Thermal Expansion
(x 10-6) |
||
|
°C
|
°F
|
/°C
|
/°F
|
|
21-93
|
70-200
|
16.5
|
9.17
|
|
21-204
|
70-400
|
16.8
|
9.35
|
|
21-316
|
70-600
|
17.0
|
9.47
|
|
21-427
|
70-800
|
17.4
|
9.64
|
|
21-538
|
70-1000
|
17.6
|
9.78
|
|
21-649
|
70-1200
|
17.8
|
9.88
|
|
21-760
|
70-1400
|
18.6
|
10.32
|
Thermal
Conductivity of A-286
The thermal conductivity at various temperatures for A-286 are given
in the following table.
|
Temperature Range
|
Thermal conductivity
|
||
|
°C
|
°F
|
W/m.K
|
Btu.ft/ft2.hr.°F
|
|
150
|
302
|
15.1
|
8.7
|
|
300
|
572
|
17.8
|
10.3
|
|
500
|
932
|
21.8
|
12.6
|
|
600
|
1112
|
23.9
|
13.8
|
Specific Heat of A-286 is 420 J/kg.K or 0.10
Btu/lb/°F
The magnetic permeability of A-286
is:
Solution treated - 1.010
Solution treated and aged - 1.007
Electrical
Resistsivity of A-286
Electrical resistivity values as a function of temperature for A-286 are given
in the following table.
|
Temperature Range
|
micro ohm-cm
|
|
|
°C
|
°F
|
|
|
25
|
77
|
91.0
|
|
540
|
1004
|
115.6
|
|
650
|
1202
|
118.8
|
|
730
|
1346
|
120.1
|
|
815
|
1499
|
122.4
|
Mechanical
Properties
Type A-286
alloy is formed most easily in the solution treated condition. Typical room
temperature tensile properties of A-286
solution treated at 1800°F (982°C) are shown below.
|
Yield Strength
|
Ultimate Tensile Strength
|
Elongation
|
|
40,000 psi
|
90,000 psi
|
40%
|
|
275 MPa
|
620 MPa
|
Cobalt
base superalloys
The use cobalt as
a constituent of metal alloys systems(not including magnets) is based on its
ability to impart high temperature strength, though this is a single phrase
which describes what can be a complex effect and in the Co/Fe/V systems to provide
controlled expansion.
The
role of cobalt is not completely understood but it certainly increases the
useful temperature range of nickel-based alloys. Y’
also occurs as Y’’ which has a body centred tetragonal
structure (i.e. two cubes stacked). Cobalt is thought to raise the melting
point of this phase thus enhancing high temperature strength.
Unalloyed cobalt has a hexagonal close-packed matrix at
temperatures below 780°F that then transforms to an FCC structure at
higher temperatures; however, nickel
alloying additions are used to stabilize the FCC austenitic structure between
room temperature and the melting point. As a class, cobalt based super-alloys are much simpler than the nickel based
alloys. Cast cobalt alloys contain about 50–60% cobalt, 20–30% chromium, 5–10%
tungsten and 0.1–1.0% car-bon. Wrought alloys contain about 40% cobalt
and high nickel contents (∼20%)
For increased workability.
Unfortunately, no precipitates that result in a large strength increase have been found for
the cobalt based alloys, and therefore they must
rely on a combination of solid solution and carbide strengthening, which limits their use in many applications. However, a
fine dispersion of carbides contributes significantly to the strength of these
alloys. In general, there are three
main classes of carbides: MC, M23C6 and M6C. In MC carbides, M stands for the reactive metals titanium, tantalum,
niobium and zirconium. In the M23C6 carbides, M is mostly chromium but can also be
molybdenum or tungsten. When the
molybdenum or tungsten content exceeds about 5 atomic percent,M6C carbides can form. The cobalt alloys display
good stress rupture properties at temperatures higher than 1830°F but cannot compete with the nickel based alloys for highly stressed parts, so are used for low stress long lived static parts. They also have superior hot
corrosion resistance at elevated temperature, probably due to their quite high
chromium contents. Examples of important cobalt based alloys are the wrought
alloys Haynes 25 (L605) and Haynes 188 and the cast alloy X-40.
The cobalt-base superalloys have their origins in the
Stellite® alloys patented in the early 1900’s by Elwood Haynes.
Although in terms of properties the (Y’) hardened nickel-based alloys have taken the lion’s share of the superalloy market, cast and wrought cobalt alloys continue to be used because:
Although in terms of properties the (Y’) hardened nickel-based alloys have taken the lion’s share of the superalloy market, cast and wrought cobalt alloys continue to be used because:
- Cobalt alloys have higher melting points than nickel (or
iron) alloys. This gives them the ability to absorb stress to a higher
absolute temperature.
- Cobalt alloys give superior hot corrosion resistance to gas
turbine atmospheres, this is due to their high chromium content.
- Cobalt alloys show superior thermal fatigue resistance and
weldability over nickel alloys.
Composition and
Structure
Cobalt alloys are termed
austenitic in that the high temperature “Face Centred Cubic” phase is
stabilised at room temperature. They are hardened by carbide
precipitation, thus carbon content is critical. Chromium provides hot corrosion
resistance and other refractory metals are added to give solid solution
strengthening – tungsten and molybdenum – and carbide formation – tantalum,
niobium, zirconium, hafnium.
Processing
is of course vital and whilst the above metals are helpful, others such as
dissolved oxygen are not. Vacuum melting is therefore becoming the norm to give
close alloy control. It is also critical that the specified compositions are
adhered to, as excess of the soluble metals, W, Mo, Cr, will tend to form
unwanted and deleterious phases similar to the nickel alloys s and Laves (Co3Ti–tetragonal close packed TCP phases).
Powder metallurgical alloys, giving a finer carbide dispersion and smaller grain size have superior properties to cast alloys. Further process development by hot isostatic pressing (HIP) has even further improved the properties by removal of possible failure sites.
Compared to nickel alloys, the stress rupture curve for cobalt alloys is flatter and shows lower strength up to about 930°C. The greater stability of the carbides, which provide strengthening of cobalt alloys, then asserts itself. This factor is the primary reason cobalt alloys are used in the lower stress, higher temperature stationary vanes for gas turbines.
Casting is important for cobalt-based alloys and directionally solidified alloys (DS) have led to increased rupture strength and thermal fatigue resistance. Even further improvements in strength and temperature resistance have been achieved by the development of single crystal alloys. Both these trends have allowed the development of higher thrust jet engines which operate at even higher temperatures.
Powder metallurgical alloys, giving a finer carbide dispersion and smaller grain size have superior properties to cast alloys. Further process development by hot isostatic pressing (HIP) has even further improved the properties by removal of possible failure sites.
Compared to nickel alloys, the stress rupture curve for cobalt alloys is flatter and shows lower strength up to about 930°C. The greater stability of the carbides, which provide strengthening of cobalt alloys, then asserts itself. This factor is the primary reason cobalt alloys are used in the lower stress, higher temperature stationary vanes for gas turbines.
Casting is important for cobalt-based alloys and directionally solidified alloys (DS) have led to increased rupture strength and thermal fatigue resistance. Even further improvements in strength and temperature resistance have been achieved by the development of single crystal alloys. Both these trends have allowed the development of higher thrust jet engines which operate at even higher temperatures.
Typical cobalt based
superalloys are HS-21, Hynes, HIperco and stellite
Satellite
Stellite
alloy is a range of cobalt-chromium
alloys
designed for wear
resistance. It may also contain tungsten
or molybdenum
and a small but important amount of carbon.
It is a trademarked name
of the Deloro Stellite Company and was invented by Elwood
Haynes. There are a large number of stellite alloys
composed of various amounts of cobalt,
nickel,
iron,
aluminium,
boron,
carbon,
chromium,
manganese,
molybdenum,
phosphorus,
sulfur,
silicon,
and titanium,
in various proportions, most alloys containing four to six of these elements.
Typical properties
Stellite alloy is a completely non-magnetic and corrosion-resistant cobalt alloy. There are a number of
Stellite alloys, with various compositions optimised for different uses.
Information is available from the manufacturer, Deloro Stellite, outlining the
composition of a number of Stellite alloys and their intended applications. The
alloy currently most suited for cutting tools, for example, is Stellite 100, because this alloy is quite hard,
maintains a good cutting edge even at high temperature, and resists hardening and annealing due to heat. Other alloys are formulated to maximize combinations
of wear resistance, corrosion resistance, or ability to withstand extreme
temperatures.
Stellite alloys display astounding hardness and toughness, and
are also usually very resistant to corrosion. Stellite alloys are so hard that
they are very difficult to machine, and anything made from them is, as a
result, very expensive. Typically, a Stellite part is precisely cast so
that only minimal machining is necessary. Stellite is more often machined by grinding, rather than by cutting. Stellite
alloys also tend to have extremely high melting points due to the cobalt and chromium
content.
Applications
Typical applications include saw teeth, hard facing, and acid-resistant machine parts. Stellite was a major improvement in the
production of poppet valves and valve seats for
the valves, particularly exhaust valves, of internal combustion engines. By reducing their erosion from hot gases, the interval between
maintenance and re-grinding of their seats was dramatically lengthened.
Stellite has also been used in the manufacture of turning tools for lathes. With the introduction and
improvements in tipped tools it
is not used as often, but it was found to have superior cutting properties
compared to the early carbon steel tools and even some high speed steel tools, especially against difficult
materials such as stainless steel. Care was needed in grinding the blanks and these were marked at
one end to show the correct orientation, without which the cutting edge could
chip prematurely.
While Stellite remains the material of
choice for certain internal parts in industrial process valves (valve seat
hardfacing), its use has been discouraged in nuclear power plants. In piping that can communicate with
the reactor, tiny amounts of Stellite would be released into the process fluid
and eventually enter the reactor. There the cobalt would be activated by the neutron flux in the reactor and become cobalt-60, a radioisotope with a five year half life that releases very energetic gamma rays. While not a hazard to the general
public, about a third to a half of nuclear worker exposures could be traced to
the use of Stellite and to trace amounts of cobalt in stainless steels.
