ASM Metals HandBook Vol. 8 - Mechanical Testing and Evaluation
Подождите немного. Документ загружается.
Type 304H
(a) In the fired section, tube surface temperatures are typically 20–30 °C (35–55 °F) higher than the tube
midwall temperature. In a typical U.S. utility boiler, the maximum metal surface temperature is approximately
625 °C (1155 °F).
Table 2 Suggested maximum temperatures in petrochemical operations for continuous service based on
creep or rupture data
Maximum temperature based on creep
rate
Maximum temperature based on
rupture
Material
°C °F °C
°F
Carbon steel 450 850 540
1000
C-0.5 Mo steel 510 950 595
1100
2 Cr-1Mo steel
540 1000 650
1200
Type 304 stainless steel 595 1100 815
1500
Alloy C-276 nickel-base
alloy
650 1200 1040 1900
The purpose of this article is to review the basic equipment and methods for creep and creep-rupture testing.
The article begins with a brief review of steady state creep behavior with subsequent sections on testing
equipment, constant-load testing, and constant-stress testing. However, another engineering aspect of creep
behavior is stress relaxation. Stress relaxation is the reduction of applied or residual stresses as a function of
time when a stressed material is held at constant strain (i.e., held at a constant length without further
deformation). Stress relaxation occurs over time, because the elastic strain within a stressed material is
transformed into permanent (i.e., plastic) strain by creep deformation. This transformation, in effect, causes a
reduction in elastic strain, which in turn, results in a reduction of the applied stress over time. For example, a
common practical example of stress relaxation is when the tightness of high-temperature bolting undergoes
progressive stress relaxation at a rate that depends on the creep strength of the material.
Stress relaxation testing principally has been used to measure and predict the relaxation of stress as a function
of time and temperature. However, stress relaxation testing can also be considered as a type of creep test.
Traditional creep-test methods involve the measurement of creep strain as a function of time with a fixed load
(or stress) at constant temperature. Stress relaxation is a test of the same behavior but under the opposite set of
test conditions, where the reduction of stress is measured with the total strain held constant at a given
temperature. Stress relaxation testing is thus a complementary creep-test method that, in a short time, can
generate creep-rate data as a function of stress covering five or more orders of magnitude in creep rate (see the
section “Accelerated Testing” in the article “Stress Relaxation Testing” in this Volume).
Footnote
*
The section “Creep Properties” was adapted from Ref 1 with a
dditional content by R.W. Hayes, Metals
Technology, Inc. The remaining sections of this article were adapted from “Creep, Stress-Rupture, and Stress-
Relaxation Testing” in Volume 8 of the 9th Edition Metals Handbook. The section “Constant-Stress Testing”
was updated by N.L. Carroll, Applied Test Systems Inc.
Reference cited in this section
1. R. Viswanathan, Damage Mechanisms and Life Assessment of High-Temperature Components, ASM
International, 1989, p 59–69
Creep and Creep-Rupture Testing*
Creep Properties
Creep properties are generally determined by means of a test in which a constant uniaxial load or stress is
applied to the specimen and the resulting creep strain is recorded as a function of time. Typical shapes of creep
curves are shown in Fig. 1. After the instantaneous strain, ε
0
a decelerating strain-rate stage (primary creep)
leads to a steady minimum creep rate (or secondary creep strain rate, ), which is finally followed by an
accelerating stage (tertiary creep) that ends in fracture at a rupture time, t
r
. The strain at rupture, ε
r
, represents
the rupture ductility.
Fig. 1 Schematic illustration of creep-curve shapes. Source: Ref 1
The shape of the creep curve is determined by several competing reactions, including:
1. Strain hardening
2. Softening processes such as recovery, recrystallization, strain softening, and precipitate overaging
3. Damaging processes, such as cavitation and cracking, and specimen necking.
Of these factors, strain hardening tends to decrease the creep rate, whereas the other factors tend to increase the
creep rate. The balance among these factors determines the shape of the creep curve. During primary creep, the
decreasing slope of the creep curve is attributed to strain hardening. Secondary-stage creep is explained in
terms of a balance between strain hardening and the softening and damaging processes, resulting in a nearly
constant creep rate. The tertiary stage marks the onset of internal- or external-damage processes (item 3 in the
preceding numbered list), which result in a decrease in the resistance to load or a significant increase in the net
section stress. Coupled with the softening processes (item 2), the balance achieved in stage 2 is now offset, and
a rapidly increasing tertiary stage of creep is reached.
Footnote
*
The section “Creep Properties” was adapted from Ref 1
with additional content by R.W. Hayes, Metals
Technology, Inc. The remaining sections of this article were adapted from “Creep, Stress-Rupture, and Stress-
Relaxation Testing” in Volume 8 of the 9th Edition Metals Handbook. The section “Constant-Stress Testing”
was updated by N.L. Carroll, Applied Test Systems Inc.
Reference cited in this section
1. R. Viswanathan, Damage Mechanisms and Life Assessment of High-Temperature Components, ASM
International, 1989, p 59–69
Creep and Creep-Rupture Testing*
Stress and Temperature Dependence
Of all the parameters pertaining to the creep curve, the most important for engineering applications are the
secondary creep, or minimum strain rate (
min
), and the time to rupture (t
r
). Specifically, their dependence on
temperature and applied stress are of utmost interest to the designer. This dependence varies with the applicable
creep mechanism. A variety of mechanisms and equations have been proposed in the literature and have been
reviewed elsewhere (Ref 2, 3, 4). Fortunately, all these mechanisms can be fitted into two basic categories:
diffusional creep and dislocation creep.
In diffusional creep, diffusion of single atoms or ions either by bulk transport (Nabarro-Herring creep) or by
grain-boundary transport (Coble creep) leads to a Newtonian viscous type of flow. The flow of atoms is
envisaged to occur from regions of local compressive stress toward regions of local tensile stress, balanced by a
counterflow of vacancies in the opposite direction. No motion of dislocations is envisaged. In this form of
creep, the steady state creep rates are postulated to vary linearly with stress (i.e., ∝ σ. At low stresses,
diffusional creep is seen only at very high temperatures above 0.6 T
m
, where T
m
is the melting point of the
material. This high-temperature regime is not generally of engineering significance for design of components.
For structural applications and high-temperature components, the pertinent temperature regime is about 0.4 to
0.6 T
m
(Ref 5). Most engineering applications for high-temperature components fall within this intermediate
temperature regime of 0.4 to 0.6 T
m
.
Creep deformation in the “intermediate” temperature regime can be associated with a combination of
mechanisms from the low-temperature regime (T < 0.4 T
m
) and high-temperature regime (T > 0.6 T
m
). At
temperatures below 0.4 T
m
, deformation is primarily controlled by dislocation glide. Recovery processes within
the low-temperature regime involve the cross slip of dislocations and are therefore considered as dynamic
recovery processes. Within the low-temperature regime, diffusional processes are considered negligible and do
not contribute to the deformation process. As temperatures are increased above 0.4 T
m
, dislocations move on
their glide planes just as at low temperatures, although more systems may become operative due to the
increased temperature. The most important difference, however, is that now dislocations acquire a new degree
of freedom, the ability to climb from one slip system to another by diffusional processes, a static recovery
process. This contribution of diffusion to creep deformation has been discussed extensively with much of our
current understanding coming from the work of Sherby (Ref 5). Sherby and Miller (Ref 6) have also compiled a
large body of data that clearly suggest a one-to-one relationship between the activation energy for lattice self-
diffusion and the activation energy for creep deformation.
Power Law Model of Steady State Creep Rates. In the intermediate temperature regime (0.4 T
m
< T < 0.6 T
m
),
the creep rate varies nonlinearly with stress, as either a power function or an exponential function of stress, in
contrast to the linear stress dependence of diffusional creep. At stresses and temperatures of interest to the
engineer, the following behavior proposed by Norton (Ref 7) and Bailey (Ref 8) is generally obeyed:
= Aσ
n
(Eq 1)
where A and n are stress-independent constants. In addition, because one contribution to creep is a thermally
activated diffusion process, its temperature sensitivity would be expected to obey an Arrhenius-type expression,
with a characteristic activation energy Q for the rate-controlling mechanism. Equation 1 can, therefore, be
rewritten as (Ref 9):
(Eq 2)
where A and n are constants (for a given condition), and R is the universal gas constant.
Equation 2 is a power-law relation that provides a basis for modeling creep rates as a function of temperature
and stress. However, A, n, and Q are parameters that depend on the particular material condition, stress level,
and temperature. The parameter A, for example, includes microstructural factors that influence the modes of
creep deformation. Therefore, the parameter A must be determined from data in order to calculate absolute
values of creep rates (see “Deformation-Mechanism Maps” in this article for a discussion on how the power-
law regime may vary for a material).
Creep deformation results also indicate that the values for n and Q in Eq 2 are both variable with respect to
stress and temperature. An example of the change in the value of n is shown in Fig. 2 for a normalized-and-
tempered 1 Cr- Mo steel (Ref 10). A distinct break in the curve is evident, with n = 4 at low stresses and n =
10 at higher stresses. The breaks in the curves occurred at stresses at which the fracture mode changed from
intergranular (I) to transgranular (T) at high stresses.
Fig. 2 Variation of minimum creep rate with stress for a normalized-and-tempered 1 Cr- Mo steel. T
and I denote transgranular and intergranular failure, respectively. Source: Ref 10
Values of n ranging from n = 1 at low stresses to n = 14 at high stresses have been reported (Ref 10 11 12 13 14
15 16), and Table 3 is a sample of studies selected to illustrate this point. Although many investigators report a
distinct break in the curve, others view the value of n as continuously changing with stress and temperature.
Table 3 Examples of reported values of stress exponent and activation energy for creep of steels
Temperature
(T)
Coefficients in
the low-stress
region
Coefficients in
the high-stress
region
System
°C °F n Q, kJ/mole n Q,
kJ/mole
Interpretation of
coefficients
Reference
1 Cr-
Mo steel
510–
620
950–
1150
4 400 10 625 Grain boundary sliding at
low stresses and matrix
deformation at high stresses
10
2 Cr-
1Mo steel
565 1050 2.5 … 12 … Deformation governed by
matrix deformation
11
1Cr- Mo
steel,
heat-
affected
zone
550–
605
1020–
1120
3 300 6 300 Diffusive mechanism at low
stresses and dislocation
mechanism at high stresses
12
1Cr- Mo
steel,
base
metal
550–
605
1020–
1120
5.6 … 5.6 … …
13
Cr-Mo-
V-steel
550–
600
1020–
1110
4.9 326 14.3
503 Dislocation climb over
particles at low stresses and
bowing between particles at
high stresses
14 (p
230)
20Cr-
25Ni-Nb
steel
750 1380 3–
4.7
465–532 8–
12
440–494
Metallographic measurement
showed that
gb
= Aσ
m
,
where m = 3.4. Transition
from low-σ to high-σ
behavior was attributed to
change from grain-boundary
sliding to matrix
deformation.
14 (p
106)
20Cr-
25Ni-Nb
steel
750 1380 2–5 250–390;
average, 320
… … n ~ corresponds to boundary
sliding; n > 3 corresponds to
matrix deformation. From
scratch displacements it was
found that
gbs
= A
3.7
and
Q
gbs
= 385 kJ/mole.
14 (p
122)
20Cr-
25Ni-Nb
steel
700–
750
1290–
1380
n varied from 6 to
8.4 with increasing
T. Q varied with σ.
At σ = 79 MPa, Q
= 678 kJ/mole.
… … Values of n and Q are only
“apparent” unless a back
stress due to NbC
precipitate is considered.
14 (p 86)
To account for these changes, it has been suggested (Ref 11, 13, 14, 17, 18) that the effective stress changes
with test conditions due to changes in an “internal back stress σ
0
” and that the stress term in Eq 2 should be
modified to (σ - σ
0
)
n
. The internal back stress represents a resisting force and is postulated to arise from a
variety of microstructural factors such as dislocation configurations, precipitate dispersion, and solid-solution
effects.
While discussions continue regarding the natures of n and Q and the reasons for their variations, industrial
practice has continued to ignore these controversies and to use a simple power law (Eq 2) with discretely
chosen values of n and Q. Because variations in n and Q are generally interrelated and self-compensating, no
major discrepancies in the end results have been noted in the low-stress region, as indicated in Table 3. At
higher stress levels, the power-law model of creep strain rates begins to break down, as discussed later in the
section “Power-Law Breakdown.”
Deformation-Mechanism Maps. With a knowledge of the stress and temperature dependence of the creep rate
for each mechanism, it is possible to construct plots showing the regimes for the various mechanisms in the
stress/temperature space. These plots, which are usually called deformation-mechanism maps or Ashby maps,
are constructed by plotting the shear stress (σ
s
) normalized by the shear modulus (G), against the homologous
temperature. Ashby maps for a turbine-blade alloy (MAR-M 200) are shown in Fig. 3 (Ref 1). As discussed
previously, the diffusional creep mechanism operates at low stresses and high temperatures, whereas at
intermediate stresses, dislocation creep or power-law creep is operative. By comparing Fig. 3(a) and (b), it can
be seen that increasing the grain size from 100 μm (3.9 mils) to 1 cm (0.39 in.) expands the power-law creep
regime and appreciably decreases the creep rate of the turbine-blade material. These maps are thus very useful
in providing insight into the operative creep mechanisms.
Fig. 3 Ashby deformation maps for MAR-M 200. A turbine blade will deform rapidly by boundary
diffusion at a grain size of (a) 100 μm but not a grain size of (b) 1 cm
The two maps in Fig. 3 for the same material with different grain sizes also illustrate the wide variation in creep
rates at a given combination of stress and temperature. This wide variation is reflected by the determination of
parameter A (in Eq 2), which must be determined experimentally for calculation of creep rates.
Power-Law Breakdown. Steady state creep rate at high temperatures can be described well by a power law (Eq
2), where the stress exponent (n) ranges from 4 to 5 for face-centered cubic (fcc) metals (Fig. 4) and about 3 for
body-centered cubic (bcc) metals (Fig. 5), and where the activation energy (Q) for creep is equal to that for
lattice self-diffusion. The data presented in Fig. 4 and 5 (Ref 19) indicate that this holds for a wide variety of
pure metals for normalized stress values below about σ
s
/G = 10
-3
.
Fig. 4 Creep data for several fcc metals plotted as a function of normalized shear stress (σ
s
/G) compared
with a power-law stress exponent of n = 4. Because the activation for creep (Q in Eq 2) is the same as that
for diffusion, the term exp (-Q/RT) in Eq 2 is replaced here by the diffusion coefficient, D. b = Burgers
vector. Source: Ref 5 with data largely from Ref 19
Fig. 5 Creep data for several bcc metals plotted as a function of normalized shear stress (σ
s
/G) compared
with a power-law stress exponent of n = 3. Source: Ref 5 with data largely from Ref 19
At normalized stresses above (σ
s
/G) = 10
-3
, creep rates begin to increase more strongly with stress, and the
power-law model breaks down. This failure of the power law at higher stress levels is known in the literature as
the power-law breakdown. In this high-stress region, an exponential or sink function appears to be more
applicable. For example, an exponential relationship, although not generally used, has been proposed (Ref 9) to
explain the behavior at very high stresses, as follows:
= A exp (C
7
σ)
(Eq 3)
where A and C
7
are stress-independent constants. A sink function has also been used in the high-stress region
(see Fig. 5 in the article “Creep Deformation of Metals, Polymers, Ceramics, and Composites” in this Volume).
In addition to the increased effect of stress on steady state creep rates, the activation energies for creep within
the high-stress regime also tend to decrease to values well below those for lattice self-diffusion. Robinson and
Sherby (Ref 20) were among the first to suggest increasing stress exponents and decreasing activation energies
at the stress levels associated with power-law breakdown. This has led to the concept of effective diffusivity.
The effective diffusion coefficient is higher than the lattice diffusion coefficient, and strong evidence in favor
of the effective diffusion concept can be found in the work of Luthy et al. (Ref 21). It was found that the steady
state strain rates within the high stress-low temperature regime could be correlated with those obtained within
the high-temperature regime by using an effective diffusivity to normalize the steady state strain rate. However,
the concept of an effective diffusivity does not fully account for power-law breakdown. Spingarn et al. (Ref 22)
have pointed out that the stress exponent and the activation energy should change together within the power-
law breakdown regime. In analyzing the creep behavior of several metals within the intermediate regime,
Springarn et al. (Ref 22) found that for some metals, the activation energy decreased at lower temperatures, but
the stress exponent remains constant. In other cases, the stress exponent is found to increase with decreasing
temperature, but the activation energy remains constant. In the case of nickel, the activation energy decreases
while the stress exponent increases as expected; however, the transition temperatures and stress levels are not in
agreement with the predictions of the effective diffusivity. Perhaps the most limiting factor regarding the
effective diffusivity concept is that it does not provide a natural explanation of the increased effect of stress in
the breakdown regime.
Composite Model and Constant-Structure Creep. At sufficiently high temperatures where creep is controlled by
lattice self-diffusion steady state creep can be described quite well using a power-law approach. However, at
higher stress levels (or lower temperatures), the power-law description begins to break down. As just described,
an effective diffusion coefficient into the power-law equation provides an explanation for the increasing stress
exponents and decreasing activation energies obtained within the high-stress/low-temperature region, but in
many cases, it is not consistent with experiment.
A composite model of plastic flow, which involves diffusion-controlled recovery at high temperatures and
cross-slip-controlled recovery at low temperatures, has been developed to unify the high- and low-temperature
regimes (Ref 1, 23). In the framework of the composite model, power-law breakdown is considered as a natural
transition from diffusion-controlled recovery at high temperatures (low stresses) to cross-slip-controlled
recovery at lower temperatures (higher stresses).
The composite model is also the foundation upon which the concept of constant structure creep is derived.
Constant structure creep, as it fits within the framework of the composite model, describes the kinetics of
dislocation glide, multiplication, and annihilation within the cell interiors. Gibeling and coworkers (Ref 24, 25,
26, 27, 28) have argued that the power-law approach does not provide an adequate description of creep under
constant structure condition.
Instead, it was suggested that the creep rate under constant structure conditions (
cs
) can best be described by an
exponential kinetic law of the form derived by Kocks et al. (Ref 25) for obstacle-controlled dislocation glide.
Stress and Temperature Dependence of Time to Rupture. The behavior of t
r
with respect to σ and T is similar to
that of , with the differences being that the signs are reversed for the stress exponent and the activation energy.
The parameter A in Eq 2 also has slightly altered values. This behavior occurs when and t
r
are inversely related
through Monkman-Grant-type relationships (i.e., t
r
constant).
Footnote
*
The section “Creep Properties” was adapted from Ref 1 with additional cont
ent by R.W. Hayes, Metals
Technology, Inc. The remaining sections of this article were adapted from “Creep, Stress-Rupture, and Stress-
Relaxation Testing” in Volume 8 of the 9th Edition Metals Handbook. The section “Constant-Stress Testing”
was updated by N.L. Carroll, Applied Test Systems Inc.
References cited in this section
1. R. Viswanathan, Damage Mechanisms and Life Assessment of High-Temperature Components, ASM
International, 1989, p 59–69
2. F. Garofalo, Fundamentals of Creep and Creep Rupture in Metals, M. Fine, J. Weertman, and J.R.
Weertman, Ed., MacMillan Series in Materials Science, The MacMillan Co., New York, 1965
3. S.K. Mitra and D. McLean, Work Hardening and Recovery in Creep, Proc. R. Soc., Vol 295, 1966, p
288–299
4. G. Bernasconi and G. Piatti, Ed., Creep of Engineering Materials and Structures, Applied Science
Publishers, London, 1979
5. O.D. Sherby, Acta Metallurgica, Vol 10, 1962, p 135–147
6. O.D. Sherby and A.K. Miller, ASME Journal of Materials and Technology, Vol 101, 1979, p 387
7. F.H. Norton, The Creep of Steel at High Temperatures, McGraw-Hill, 1929
8. R.W. Bailey, Creep of Steel under Simple and Compound Stress, Engineering, Vol 121, 1930, p 265
9. O.D. Sherby and P.M. Burke, Mechanical Behavior of Crystalline Solids at Elevated Temperatures,
Prog. Mater. Sci., Vol 13, 1967, p 325
10. R. Viswanathan, The Effect of Stress and Temperature on the Creep and Rupture Behavior of a 1.25 pct.
Chromium -0.5 pct. Molybdenum Steel, Metall. Trans. A, Vol 8A, 1977, p 877–883
11. B.J. Cane, “The Process Controlling Creep and Creep Fracture of 2 Cr-1Mo Steel,” CEGB Report
RD/LR 1979, Central Electricity Generating Board Research Laboratories, Leatherhead, England, 1979
12. M.S. Shammas, “Estimating the Remaining Life of Boiler Pressure Parts,” EPRI Final Report on RP
2253-1, Vol 4, Electric Power Research Institute, Palo Alto CA, 1987
13. M.C. Askins et al., EPRI Final Report on RP 2253-1, Vol 5, Electric Power Research Institute, Palo
Alto, CA, 1987
14. I.R. McLauchlin, in Creep Strength in Steel and High Temperature Alloys, The Metals Society, London,
1974, p 8, 71, 86, 106, 122, 217, 230
15. K.E. Amin and J.E. Dorn, Acta Metall., Vol 17, 1969, p 1429
16. V. Foldyna, A. Jakobova, T. Prnka, and J. Sabotka, Creep Strength in Steels and High Temperature
Alloys, The Metals Society, London, 1974, p 130
17. K.R. Williams and B. Wilshire, Metall. Sci. J., Vol 7, 1973, p 176
18. C.N. Aneceuist, R. Gasca-Neri, and W.D. Nix, A Phenomenological Theory of Steady State Creep
Based on Average Internal and Effective Stress, Acta Metall., Vol 18, June 1970, p 663–671
19. A.K. Mukherjee, J.E. Bird, and J.E. Dorn, ASM Trans. Quart., Vol 62, 1969, p 155
20. S.L. Robinson and O.D. Sherby, Acta Metall., Vol 17, 1969, p 109
21. H. Luthy, A.K. Miller, and O.D. Sherby, Acta Metall., Vol 28, 1980, p 169
22. J.R. Spingarn, D.M. Barnett, and W.D. Nix, Acta Metall., Vol 27, 1979, p 1549
23. W.D. Nix, J.C. Gibeling, and D.A. Hughes, Metall. Trans. A., Vol 16, 1985, p 2215
24. M. Biberger and J.C. Gibeling, Acta Metall. Mater., Vol 43, 1995, p 3247
25. U.F. Kocks, A.S. Argon, and M.F. Ashby, Thermodynamics and Kinetics of Slip, Pergamon Press,
Oxford, 1975, p 141
26. G.S. Nakayama and J.C. Gibeling, Acta Metall., Vol 38, 1990, p 2023
27. G.S. Nakayama and J.C. Gibeling, Scr. Metall. Mater., Vol 24, 1990, p 2031
28. G.S. Nakayama and J.C. Gibeling, Modeling the Deformation of Crystalline Solids, T.C. Lowe et al.,
Ed., TMS, Warrendale, PA, 1991, p 589
Creep and Creep-Rupture Testing*
Extrapolation Techniques
Because many high-temperature components are designed to operate over long periods of time (sometimes in
excess of 100,000 h), extrapolation of laboratory creep and rupture data to actual service conditions is
unavoidable. Even if long-time data are available for selected heats of material, heat-to-heat variations in
properties make it necessary to estimate the long-time behavior for other heats. Greater difficulty is encountered
in estimating the remaining creep lives of in-service components, where decisions have to be made based on
very short-time laboratory tests (usually less than 1000 h). The need for extrapolation techniques that permit
estimation of the long-term creep and rupture strengths of materials based on short-duration tests is thus a very
real and important one in design, quality control, and plant evaluation.
Basically, parametric techniques incorporate time stress and temperature test data into a single expression.
When test data recorded over adequate times and at temperatures above the service temperature are
incorporated into a single “master curve,” the stress for the service-temperature conditions can be read directly
from the master curve.
Extensive reviews of parametric techniques have been given elsewhere (Ref 29 and 30). Their historical
evolution up to the present time has been reviewed by Manson and Ensign (Ref 31). More detailed information
on extrapolation is also provided in the article “Assessment and Use of Creep-Rupture Properties” in this
Volume.
Footnote
*
The section “Creep Properties” was adapted from Ref 1
with additional content by R.W. Hayes, Metals
Technology, Inc. The remaining sections of this article were adapted from “Creep, Stress-Rupture, and Stress-
Relaxation Testing” in Volume 8 of the 9th Edition Metals Handbook. The section “Constant-Stress Testing”
was updated by N.L. Carroll, Applied Test Systems Inc.
References cited in this section
29. J.B. Conway, Stress Rupture Parameters: Origin, Calculations and Use, Gordon and Breach, 1969
30. S.S. Manson, Time-Temperature Parameters—A Re-Evaluation of Some New Approaches, ASM
Publication D-8-100, American Society for Metals, 1968, p 1–115
31. S.S. Manson and C.R. Ensign, A Quarter Century of Progress in the Development of Correlation and
Extrapolation Methods for Creep Rupture Data, ASME J. Eng. Mater. Technol., Vol 101, 1979, p 317–
325
Creep and Creep-Rupture Testing*
Methods and Equipment
Determination of creep characteristics of metals at high temperatures requires the use of a loading device or test
stand, an electric furnace with suitable temperature control, and an extensometer. Equipment discussed in this
section is for uniaxially loaded specimens in tension. More information on creep and stress-rupture testing
equipment and methods can be found in the following standard: