HIGH CYCLE FATIGUE STUDIES OF CARBURIZED NICKEL-BASE SUPERALLOYS AND STAINLESS STEELS
by
YINDONG GE
Submitted in partial fulfillment of the requirements For the degree of Master of Science
Thesis Advisor: Dr. Arthur H. Heuer
Department of Materials Science and Engineering CASE WESTERN RESERVE UNIVERSITY
August, 2009
CASE WESTERN RESERVE UNIVERSITY
SCHOOL OF GRADUATE STUDIES
We hereby approve the thesis/dissertation of
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Dedication
This thesis is dedicated to my dear parents Mr. Yaorong Ge and Ms. Suzhen Zhang. Table of Contents
List of Figures………………………………………………………………………..iii
List of Tables…………………………………………………………….………...... ix
Acknowledgement………………………………………………………………...... xi
Abstract…………………………………………………………………………...…xii
1. Introduction……………………………………………………………………….1
2. Background
2.1 Methods to Improve Fatigue Life…………………………….………4
2.2 Low Temperature Carburization Technique………………….…....10
2.3 Fatigue Studies of IN718, A286 and 2205………………..…….....13
3. Materials and Methods
3.1 Materials Specifications……………………….…………………….17
3.2 Tensile Testing……………….……………………………………….33
3.3 Fatigue Testing…………………………………………….…………34
3.4 Charpy Testing……………………………………………………..…35
3.5 Hardness Testing……………………………………………….……35
3.6 Fractography…………………………………………………………35
4. Results and Discussions
4.1 IN718 Ni-base Alloy…………………………………………………36
i
4.2 A286 Fe-base Alloy…………………………………………………55
4.3 2205 Duplex Stainless Steel……………………………………….72
5. Conclusions…………………………………………………………………….89
Future Work…………………………………………………………………………91
Appendix I…………………………………………………………………...……...93
Appendix II………………………………………………………………..……...…96
Bibliography…………………………………………………………………...…..122
ii
List of Figures S.No Title Page
.
2.1 S–N curves for Bare, Cr plated, 5 and 10 μm thick DLC 6
coated 4340M steel.
2.2 Variations of numbers of cycles to failures with stresses (S–N 7
curves) of (a) the 316 L stainless steel, and (b) the Ni-based
alloy with and without the film
2.3 Fatigue damage mechanism for nitrided 42CrMo4 steel grade 9
(A) and the same+CrCN coating (B). It is possible to see a
non-metallic inclusion acting as fatigue nucleation site.
2.4 Stress–strain plot of one time carburized (C), nontreated (NT), 11
and heat-treated (HT) 316L stainless steel samples
2.5 Sample geometry used for tensile and R = -1 fatigue testing 11
2.6 S–N curves shows data for nontreated (NT) and one-time (1), 12
two-times (2), and four-times (4) carburized samples.
2.7 S-N curve shows additional data for NT electropolished 13
samples and for NT electropolished samples with and without
diamond polishing (D/P E/P and E/P, respectively) and heat-
treated (HT) samples.
2.8 S-N diagrams at 4K, 77K and 293K for IN718 alloy in the 14
study
2.9 Fatigue strengths for fine- and coarse-grain alloy A286 15
2.10 S-N curves showing the fatigue limits of the studied 2205 16
iii
3.1 IN718 microstructure reveals by etchant 3 gram CuCl2 + 30 18
ml HCl + 3 ml H2O, Swabbed for 5 s
3.2 Schematic of a carburization treatment 19
3.3 Carburized IN718 surface hardness and carbon concentration 20
(hardness profile is measured by the Swagelok Company)
3.4 XRD analysis result of non-treated and carburized IN718 21
3.5 XRD residual stress measurement result of carburized IN718 21
3.6 A286 microstructure revealed by 3 Part HCl + 1 Part HNO3 + 23
1 Part Glycerol + few Drop HF, swabbed for 30-40 s (by Dr.
Shaghi-Moshtaghin)
3.7 Carburized A286 surface hardness and carbon concentration 24
(hardness profile is measured by the Swagelok Company)
3.8 XRD analysis result of non-treated and carburized A286 25
3.9 Carburized A286 XRD residual stress measurement result 25
3.10 Metallography of carburized 2205 samples. 26
3.11 AES elements profile of carburized wrought 2205 27
3.12 Carburized 2205 surface hardness 28
3.13 Plan view SEM image of carburized 2205 (a), EBSD pattern 29
from FCC phase (b) and EBSD pattern from previous BCC
phase (c).
3.14 XRD analysis result of non-treated and carburized 2205 30
3.15 Carburized 2205 XRD residual stress measurement result, 30
austenite phase (a), the other phase (b)
3.16 Geometry of samples used for tensile and fatigue tests 33
iv
3.17 MTS 810 materials testing system machine 34
4.1 Stress-strain curves of carburized, non-treated, solution- 37
treated and heat-treated IN718 samples in the low strain
range
4.2 Engineering stress-strain curves of carburized, non-treated, 38 solution treated and heat-treated IN718 in the final fracture region. Strain obtained based up crosshead displacement. 4.3 True stress-strain curves of carburized, non-treated, solution 39
treated and heat-treated IN718 in the final fracture region.
Strain obtained based up crosshead displacement.
4.4 Hardness data of IN718 sample’s grip 40
4.5 Fractograph of non-treated IN718 sample, showing internal 41 fatigue crack initiation (550 MPa) 4.6 Fatigue results for C and NT specimens. (b) Stress-strain 44
behavior before fatigue testing for the yielded C specimen
shown in (a).
4.7 SEM images of fracture surface of carburized sample (a), (b); 45
non-treated sample (c), (d). Both the samples failed at 550
MPa stress, 50% of Y.S. (e) is the shear shape at the edge of
fractured fatigue carburized sample
4.8 SEM images of a fatigue fractured HT IN718 sample, (a) and 51
(b) lateral view, (c) plan view
4.9 Log-log plot of carburized and non-treated IN718 fatigue 53
results
4.10 Stress-strain curves of batch C carburized and non-treated 56
A286 samples in the low strain range
v
4.11 Stress-strain curves of batch C carburized, non-treated A286 57
samples
4.12 True stress-strain curves of carburized, non-treated A286 57
samples
4.13 Side view image of A286 rod samples fractured in tension, (a) 59
and (b) non-treated; (c) and (d) carburized.
4.14 The batch C carburized and non-treated A286 fatigue life. 63
4.15 Log-log plot of carburized and non-treated A286 fatigue 63
results
4.16 SEM images of fracture surface of non-treated fatigue 64
samples fractured at 380 MPa (a), (b); 510 MPa (c), (d), (e).
4.17 SEM images of fracture surface of carburized sample fatigue 67
fractured at 380 MPa (a), (b); 470 MPa (c), (d).
4.18 Stress-strain curve of carburized and non-treated 2205 73
samples in the low strain range
4.19 Stress-strain curves of carburized, non-treated 2205 samples 74
4.20 True stress-strain curves of carburized, non-treated 2205 74
samples
4.21 Tension fracture surface and neck area of 2205 samples, (a) 76
and (b) non-treated; (c) and (d) carburized
4.22 S-N curve shows the carburized and non-treated 2205 80
samples fatigue life.
4.23 Comparison S-N curves of the non-treated 2205 samples in 80
our group and L samples in Mateo’s group
4.24 Log-log plot of carburized and non-treated 2205 fatigue 79
vi
results
4.25 SEM images of fracture surface of sample fatigue fractured at 83
480 MPa: non-treated (a), (b); carburized (c), (d); (e) shows
the XEDS spectrum and chemical composition of inclusion
particle.
4.26 SEM images of fracture surface of non-treated sample fatigue 86
fractured at 440 MPa; (c) shows the XEDS spectrum of
inclusion particle.
A1 XRD lattice parameters for the (420) peak as a function of 94
sin2ψ.
A2 The engineering stress-strain curves of non-treated and 97
carburized 316L wires. a) 76 um thick wires; b) 127 um thick
wires; c) 203 um
A3 Fracture surface of non-treated 316L wires a) 76 um, b) 203 99
um; and carburized wires c) 76 um, d) 203 um.
A4 Stress-strain curves of different surface finishes carburized 102
316L samples in the low strain range.
A5 Stress-strain curves of different surface finishes carburized 103
316L samples.
A6 S-N curve shows the different surface finishes carburized 104
316L fatigue life.Tests conducted with a load ratio, R=-1.
A7 Stress-strain curve of non-treated and carburized MP98T 106
samples in low strain range.
A8 Stress-strain curves of non-treated and carburized MP98T 107 samples. A9 The fatigue testing results for non-treated and carburized 108 MP98T. Tests conducted with a load ratio, R=-1. vii
A10 Stress-strain curves for non-treated and carburized 303 110
samples in low strain range.
A11 Stress-strain curves for non-treated and carburized 303 111
samples
A12 The fatigue testing results for non-treated and carburized 303 112 stainless steel samples. Tests conducted with a load ratio, R=- 1. A13 Stress-strain curve of non-treated and carburized AL6XN 115
samples in the low strain range
A14 Stress-strain curves of non-treated and carburized AL6XN 116
samples
A15 The fatigue testing results obtained from non-treated and 117
carburized AL6XN samples.
A16 Photograph of carburized A286 flat bar 118
A17 Yield curve of carburized A286 flat tension bar 119
A18 XRD measurement of surface residual stress magnitude 120
viii
List of Tables S.No Title Page
.
2.1 Mechanical properties of alloys 4
2.2 Tensile properties in IN718 used in study 13
2.3 High-temperature tensile properties of materials used 15
2.4 Tensile tests results of the investigated 2205 16
3.1 Chemical composition of IN718, A286 and 2205 17
3.2 Estimated volume expansion, induced stress according to 31
corresponding carbides
4.1 Uniaxial tensile tests result of carburized, non-treated, 37
solution-treated IN718 superalloy
4.2 Charpy test results of non-treated and carburized IN718 42
4.3 Basquin parameters for carburized and non-treated IN718 53
samples
4.4 Uniaxial tensile tests data of carburized, non-treated A286 55
superalloy
4.5 Charpy test results of non-treated and carburized A286 61
4.6 Basquin parameters for carburized and non-treated A286 62
samples
4.7 Uniaxial tensile tests data of non-treated, carburized and heat- 72
treated A286 superalloy
4.8 Charpy test results of non-treated and carburized 2205 78
4.9 Basquin parameters for carburized and non-treated 2205 82
ix
samples
5.1 Basquin parameters for carburized and non-treated 2205, 91
IN718 and A286
A1 Uniaxial tensile tests conducted on carburized 316L stainless 101
steel samples having different surface finishes
A2 Chemical Composition of MP98T 105
A3 Uniaxial tensile tests conducted with the MP98T superalloy in 106
non-treated and carburized conditions
A4 Uniaxial tensile test results for non-treated and carburized 303 109
stainless steels
A5 Chemical Composition of AL6XN 113
A6 Uniaxial tensile test results for non-treated and carburized 114
AL6XN stainless steels
A7 Vickers hardness data obtained using 25 gram force 121
x
ACKNOWLEDGEMENTS
First of all, I wish to thank my thesis advisor, Professor Arthur H. Heuer, for his guidance and supervision throughout this study. I am also very grateful to
Professor Gary M. Michal, Professor Frank Ernst, Professor John J.
Lewandowski and Professor Harold Kahn for many helpful discussions and advice. I also hope to express my appreciation to Dr. Reza Shaghi-
Moshtaghin for his support in obtaining data in my thesis. I would like to thank
Mr. Christopher Tuma as well as my colleagues in the Materials Science and
Engineering Department, who gave me numerous amount help and support.
xi
High Cycle Fatigue Studies of Carburized
Nickel-base Superalloys and Stainless Steels
Abstract
by
YINDONG GE
Fatigue properties of non-treated and carburized Ni-base IN718, Fe-base
A286 and duplex stainless steel 2205 were investigated in fully-reversed uniaxial fatigue tests. The carburization process changes the fatigue crack initiation site of IN718 from surface-nucleated for the non-treated samples to centrally-nucleated for carburized samples. However, the fatigue lifetimes are not increased, because of defects that exist in the central areas of all the specimens functioning as weak points that served as sites for fatigue crack nucleation. The fatigue resistance of carburized A286 samples is modestly improved. Additionally, the fatigue strength coefficient and exponent of A286 are changed by carburization, which also effectively suppresses the formation of multiple crack nucleation sites at the surface. Carburized 2205 samples exhibited significant improvement in fatigue limit and fatigue lifetime compared to non-treated samples at the same cyclic stress, in spite of the existence of oxide inclusions, which act as centers for fatigue crack nucleation.
xii
1. Introduction
In materials science, high cycle fatigue is the progressive, localized, and permanent structural damage that occurs when a material is subjected to cyclic or fluctuating strains at nominal stresses that have maximum values less than (and often much less than) the static yield strength of the material.
The resulting stress is thus well below the ultimate tensile stress, or even the yield stress of the material, and yet still can cause fracture. Basically, at first, one or more small cracks initiate in the material. With cyclic loading, damage will accumulate, and the dominant cracks grow resulting in the material’s final
fracture.
It is noteworthy that fatigue failure is a major component of the total costs of
fracture and its prevention. Fatigue cracking under low-amplitude, high-cycle loading is the dominant failure process in a number of engineering applications. Moreover, an estimated 80% of these costs involve situations
where cyclic loading and fatigue are at least a contributing factor. [1]
For example, the principal cause of failure of components, from which a gas
turbine jet engine of a modern military aircraft is made, is high-cycle fatigue.
As reported, fatigue failure accounts for 49% of all component damage in jet
engines. High-cycle fatigue (HCF) is responsible for half of all the failures,
while low–cycle fatigue (LCF) and all other modes of fatigue lead to the
remainder of fatigue failures in roughly equal proportions. [2]
1
In recent year, many structural materials are required to have good mechanical properties like wear and fatigue resistance as well as excellent corrosion resistance due to the demand of a severe working environment for components. Stainless steels have an advantage of remarkable corrosion resistance. However, poor wear and fatigue properties significantly constrain the wide usage of stainless steels. IN718 as well as other superalloys have wear characteristics that are also unsatisfactory, though they have otherwise excellent properties. Thus, surface treatments improving their tribological behavior without adversely effecting corrosion resistance are developed to enhance performance and expand the field of applications of superalloys.
A variety of surface engineering techniques have been developed with the aim of making hard surfaces. When surface-modified materials are applied to load-bearing components, the fatigue properties become critical. The fatigue properties of materials after different surface treatment methods such as plasma nitriding, gaseous nitriding, carburization, TiN film deposition and duplex (combination of two methods) have been reported.
Fatigue cracks usually initiate at the surface due to defects and stress concentration at surface. For this reason, the conditions of the surface and its environment are important factors influencing fatigue behavior. Plastic deformation is essential to the nucleation and growth of fatigue cracks.
Anything that promotes plastic deformation will decrease the fatigue resistance of a material, and vice versa, fatigue resistance can be enhanced by inhibiting plastic deformation. Basically, suppression of surface crack
2 initiation can remarkably improve the fatigue properties of materials.
In this study, the Low Temperature Colossal Supersaturation (LTCSS) process, which was invented by the Swagelok Company, was employed to improve materials’ fatigue and wear resistance and surface hardness.
Compared to the conventional carburization process, the LTCSS process can generate a case containing about 10 at% or even higher carbon in solid solution and to depths about 15-20 μm depending on the material’s characteristic, while suppressing the formation of carbides, resulting in high surface hardness without reducing corrosion resistance.
A significant compressive residual stress (up to 3 GPa) is generated in the
LTCSS process. Experimental results prove that this high level of residual stress is very beneficial to the improvement of fatigue life of materials. In this study, the effect of carburization on the fatigue properties of Ni base superalloy IN718, precipitation hardened Fe-base superalloy A286 and duplex stainless steel 2205 have been studied. Different degrees of improvement in fatigue lives following carburization have been observed.
3
2. Background
2.1 Methods to Improve Fatigue Life
Nitriding has proven to be one of the most effective methods to enhance the fatigue properties of steels. Various nitriding methods have been applied to steels, and the fatigue behavior of nitrided steels has been extensively investigated.
A liquid nitirding process is reported to be applied to an AISI 4130 low alloy steel, followed by an investigation of its fatigue behavior by rotation bending fatigue tests and axial fatigue using the compact tension (CT) geometry [3].
The fatigue limit in rotating bending tests increase after nitriding (Table 2.1, σe is the fatigue limit) while the CT tests showed that the nitrided specimens had a lower resistance to crack initiation in air and in an chloride solution compared to the non-nitrided sample. It was also reported that the nitrided layer constituents, mainly FeN and CrN, are the reason for the increment in hardness and the fatigue limit. However, the porous nature of the nitrided layer lowers the crack threshold stress.
Table 2.1: Mechanical properties of alloys
4
However, not all nitriding research has found a beneficial result for the fatigue behavior of steels. Plasma nitriding is reported to generate higher surface hardness, compressive residual stresses at the surface and lower surface roughness on AISI 304 austenitic stainless steel [4].
Coating is one of the widely applied methods which are used to improve the fatigue properties of steels. Hard chromium plating is currently used to protect the surfaces of wear and fatigue-sensitive aircraft parts such as the landing gear, flap tracks, etc. [5]. Application of a Diamond-Like Carbon (DLC) coating to improve wear and fatigue properties has also been studied. DLC films were deposited by a high frequency plasma process using a mixture of hydrocarbons such as methane and inert gases such as argon at subatmospheric pressures [6]. Tests show that a DLC film has excellent adhesion to a steel surface and DLC coated 4340 M steel exhibited better wear and fatigue properties than those of chromium-plated 4340 M (Figure
2.1).
5
Figure 2.1: S–N curves for Bare, Cr plated, 5 and 10 μm thick DLC coated
4340M steel. [6]
Zr-based glass-forming metallic films have proven to be very beneficial to the fatigue life of 316L stainless steel and C-2000 Ni base alloy [7]. Results of
four-point-bending fatigue tests showed that the fatigue life and fatigue-
endurance limit of the materials are significantly enhanced. High hardness
and good ductility of the glass-forming coating, coupled with good adhesion between the film and the substrate, as well as improved surface finishes, contribute to the fatigue resistance increment.
6
Figure 2.2: Variations of numbers of cycles to failures with stresses (S–N curves) of (a) the 316 L stainless steel, and (b) C-2000 Ni-based alloy with and without the film [7]
7
Recently, a tendency to use combinations of two or even more than two
treatment methods has been evident. The influence of plasma nitriding, Ti-
DLC thin film deposition, and duplex surface treatment (nitriding +Ti-DLC deposition) on the fatigue life of 316L was investigated. Compared to the untreated samples, all treated samples had improved fatigue strength, while duplex treatments attained the best fatigue strength [8].
The resistance of low-carbon steel 20 to fatigue after a combined process of surface hardening, including laser alloying, followed by nitriding was analyzed.
And a combination of high hardness, strength, resistance to fatigue failure, and residual stress was achieved [9].
Sometimes, duplex treatments result in less improvement than one treatment.
It is reported that Cr PVD coating on 42CrMo4 steel induces a high compressive residual stress zone which increases the fatigue resistance [10].
On the contrary, a duplex-treatment, a combination of nitriding pre-treatment and Cr PVD coating, modestly increased the fatigue limit above that of uncoated nitrided samples. This phenomenon is believed due to the residual stress gradient which does not show significant differences for nitrided coated and uncoated samples. The PVD coating seems not to modify the damage mechanisms of 42CrMo4 samples during the fatigue tests. The nucleation of the fatigue damage is located in the substrate (Figure 2.3), and it is always related to a non-metallic inclusion.
8
(a)
(b)
Figure 2.3: Fatigue damage mechanism for nitrided 42CrMo4 steel grade (A) and the same+CrCN coating (B). It is possible to see a non-metallic inclusion acting as fatigue nucleation site.
Carburization is a surface hardening process where the surface carbon
9
content is increased by heating the component below its melting point with
carbonaceous matter to create a ductile core with a hard surface. A hard
carburized case layer, induced by heat treatment in a carbonaceous
atmosphere, can significantly affect commercial AISI 8620 steel’s low and high
cycle fatigue behavior [11]. The effect of the hard layer thickness in fatigue life
and fatigue strength was investigated. Fatigue properties of carburized steels
increase with the increment of case depth. A combination of transgranular and
intergranular fracture mechanisms was observed.
2.2 Low Temperature Carburization Technique
A low-temperature paraequilibrium carburization process was developed by the Swagelok company [12] and studied at Case Western Reserve University
[13,14]. Fatigue and other mechanical properties of carburized 316L austenitic stainless steel were investigated [15]. Attributed to the formation of greater than 2 GPa compressive residual stress and a three times harder case, both generated by a ‘‘colossal’’ carbon supersaturation, fatigue endurance limit increased from about 200 MPa to 325 MPa.
Figure 2.4 shows the stress–strain curves from uniaxial loading comparing the
Carburization Treated (C), Non-Treated (NT) and Heat Treated (HT) samples.
The result demonstrated that essentially no embrittlement of the sample accompanied carburization, and there was likewise no change in ductility.
10
Figure 2.4: Stress–strain plot of one time carburized (C), nontreated (NT),
and heat-treated (HT) 316L stainless steel samples
Uniaxial fully reverse tension-compression test was conducted on samples
with the geometry shown in Figure 2.5
Figure 2.5: Sample geometry used for tensile and R = -1 fatigue testing
Figure 2.6 is the fatigue test results of carburized (C), nontreated (NT), and
heat-treated (HT) 316L stainless steel samples. A dramatic enhancement of
11
the fatigue endurance limit from 200 MPa for non-treated samples to 325 MPa
for carburized ones was observed, while different numbers of treatments do
not show a significant effect (neither beneficial nor detrimental) on fatigue
limit.
Figure 2.6: S–N curves shows data for nontreated (NT) and one-time (1), two-times (2), and four-times (4) carburized samples of 316L stainless steel.
Figure 2.7 shows that different polishing techniques do not affect the fatigue lifetimes of NT samples. The heat treatment alone has almost no effect on the fatigue performance.
12
Figure 2.7: S-N curve shows additional data for NT electropolished samples and for NT electropolished samples with and without diamond polishing (D/P
E/P and E/P, respectively) and heat-treated (HT) samples of 316L stainless steel.
2.3 Fatigue Studies of IN718, A286 and 2205
Y. Ono et al. [16] conducted tensile tests and fatigue tests on IN718 at 4K,
77K and 293K. Tensile tests were carried out using cylindrical specimens 6.25 mm in diameter and 35 mm in gauge length. The fatigue specimens were an hourglass-type with a minimum diameter of 4.5 or 6 mm. Fatigue tests were performed using a sinusoidal stress ratio (R) of 0.01.
13
Table 2.2: Tensile properties in IN718 used in study Tensile Properties 4 K 77 K 293 K
0.2% yield stress (MPa) 1414 1327 1174
Ultimate tensile strength (MPa) 1822 1640 1369
Total elongation (%) 15.3 12.5 10.6
Reduction of area (%) 14.2 13.1 14.6
Figure 2.8: S-N diagrams at 4K, 77K and 293K for IN718 alloy in the study
[16]
Kobayashi et al. [17] investigated high-temperature fatigue properties of austenitic superalloys A286 in region between 102 and 107 cycles. The fatigue
14
strengths of alloys A286 in high-cycle region over 104 cycles showed a grain size effect. The initiation sites of the fracture were the crystallographic facets corresponding to the austenitic grain size.
Table 2.3: High-temperature tensile properties of materials used Materials Temperature 0.2% yield Tensile Elongation (℃) Stress (MPa) strength (MPa) (%)
Coarse-grain 600 593 837 19 alloy A286 Fine-grain 600 713 912 22 alloy A286
Figure 2.9: Fatigue strengths for fine- and coarse-grain alloy A286 [17]
Tensile and high cycle fatigue behavior of 2205 specimens machined from the plate materials were characterized by A. Mateo etc. [18] Specimens for the
15 plate were machined parallel (L), perpendicular (T) and at a 45° angle (D) to the rolling direction. A resonant testing machine was used for the HCF tests.
These were conducted under fully reversed load control.
The results from the tensile analysis are given in Table 2.4.
Table 2.4: Tensile tests results of the investigated 2205
Sample σy (MPa) σuts (MPa) Elongation (%)
T 650 840 32
L 600 770 40
D 470 690 46
Figure 2.10: S-N curves showing the fatigue limits of the studied 2205
16
3. Materials and Methods
3.1 Materials Specifications
All the IN718, A286 and 2205 samples, both non-treated and carburized, were supplied by the Naval Research Laboratory (NRL). The composition of the
IN718, A286 and 2205 are listed in Table 3.1.
Table 3.1: Chemical composition of IN718, A286 and 2205 Element at% IN718 A286 2205
C 0.03 0.08 max <0.03
Al 0.52 0.35 max -
Si 0.08 1.0 max <1
Ti 0.91 1.90-2.35 -
Cr 18.25 13.5-16.1 21-23
Fe 17.69 Balance Balance
Ni Balance 24-27.1 4.5-6.5
Nb 4.91 - -
Mo 3.00 1.0-1.5 2.5-3.5
Mn 0.1 2.0 max <2
Co 0.28 - -
Cu 0.14 - -
IN718
17
The alloy IN718 is one of the most common precipitation-hardening Ni-base
superalloys. It belongs to a class of so-called Ni-Fe superalloys, containing a
relatively high content of Fe as well as Ti and Al to provide γ’ (Ni3(Ti,Al))
precipitation. It also contains Nb which leads to precipitation of the tetragonal
Ni3Nb phase (γ’’). Alloy IN718 has a density, modulus and Poisson’s ratio of
8.22 g/cm3, 191 GPa and 0.29 respectively. [19]
Figure 3.1 is microstructure of IN718 sample carburized at 510 ℃ (taken by
Dr. Shaghi-Moshtaghin). From this image, a visible case of around 20 μm thickness is found.
20 µm
Figure 3.1: IN718 microstructure reveals by etchant 3 gram CuCl2 + 30 ml
HCl + 3 ml H2O, Swabbed for 5 s
In the carburization process, IN718 is first activated by HCl at 450 ℃ to
18 remove its chromium-rich passive film. Then the surface of the material is exposed to a carburization gas of CO, HCL and H2 at 510 ℃ for a designed period. The carburization process is shown schematically in the Figure 3.2.
Figure 3.2: Schematic of a carburization treatment [14]
The surface hardness of carburized IN718 alloy has been measured by the
Swagelok Company. Its surface carbon concentration profile was determined at Case using Auger Electron Spectroscopy (AES), and is shown in Figure
3.3.
19
Figure 3.3: Carburized IN718 surface hardness and carbon concentration
(hardness profile is measured by the Swagelok Company)
X-Ray Diffraction (XRD) was applied to both non-treated and carburized
IN718 by Dr. Kahn. His results are shown in Figure 3.4. As expected, the
carburized IN718 shows a peak shift to smaller 2Θ vaules, which indicates that carburization process expands the lattice parameter of IN718 by 1.2%.
Compressive residual stress at the carburized surface also was determined using XRD by Dr. Kahn. The method of residual stress analysis is listed in
Appendix I. Elastic constants of IN718 used in this calculation are from D.
Dye. [20] The magnitude of residual stress is 1.9 GPa, in compression as shown in Figure 3.5.
20
Figure 3.4: XRD analysis result of non-treated and carburized IN718
Figure 3.5: XRD residual stress measurement result of carburized IN718
21
A286
A286 alloy is an iron-base superalloy widely used by the aerospace industry for intermediate elevated temperature service. It is nominally a 25wt. % Ni and
15wt. % Cr austenitic alloy containing Ti, Al and other minor alloying additives.
It is mainly strengthened by an aging treatment which precipitates the ordered fcc γ’ phase, Ni3(Ti, Al), but variation in the minor alloying additives significantly influences the precipitation sequence.[21]
Type A286 alloy is a heat (up to 1300 ℉, 704 ℃) and corrosion resistant austenitic iron base material which can be age hardened to a high strength level. A286 is also useful for lower stress applications at higher temperatures.
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 ℉ (-196 ℃). The alloy may be used for moderate corrosion applications in aqueous solutions. [19]
A286 was activated in HCl prior to carburization in an atmosphere of CO. The same processing as applied to IN718 as shown in Figure 3.2 was empolyed except that the activation temperature was set at 324 ℃ and carburization temperature 450 ℃.
After etching, a carburized case with a depth of 37 μm can be easily distinguished in A286 samples. It reveals that carbon atoms diffuse
22 significantly faster in austenitic iron base alloy A286 than in Ni base alloy
IN718.
Figure 3.6: A286 microstructure revealed by 3 Part HCl + 1 Part HNO3 + 1
Part Glycerol + few Drop HF, swabbed for 30-40 s (by Dr. Shaghi-Moshtaghin)
The surface hardness of the carburized A286 alloy has been measured. The surface carbon concentration profile was determined using Auger Electron
Spectroscopy (AES). The hardness and carbon concentration profile are shown in Figure 3.7.
23
Figure 3.7: Carburized A286 surface hardness and carbon concentration
(hardness profile is measured by the Swagelok Company)
X-Ray Diffraction (XRD) was applied to both non-treated and carburized A286 by Dr. Kahn. His results are shown in Figure 3.8. The magnitude of the peak shift indicates that the carburization process expanded the lattice parameter of
A286 by 4.5%, which is a very significant expansion.
Compressive residual stress at the carburized surface was also determined using XRD by Dr. Kahn. The elastic constants for A286 used in this calculation are from H. Ledbetter. [22] The magnitude of residual stress was -3.3 GPa, as shown in Figure 3.9.
24
Figure 3.8: XRD analysis result of non-treated and carburized A286
Figure 3.9: Carburized A286 XRD residual stress measurement result
25
2205
2205 is the most widely used duplex (ferritic/austenitic) stainless steel grade.
It finds applications due to both its excellent corrosion resistance and high
strength.
Typical applications include: Chemical processing, transport and storage. Oil and gas exploration and processing equipment Marine and other high chloride environments Pulp and paper digesters, liquor tanks and paper machines
The activation and carburization temperatures for 2205 were both 380 ℃ and
the carburization time was 150 hours. Figure 3.10 shows the microstructure of
the carburized wrought 2205 sample obtained from the Swagelok Company.
The blue line indicates the path of an AES line scan.
1
2000 X 10.0 keV 10.0 µm DA019 10/20/2008
Figure 3.10: Metallography of carburized 2205 samples.
26
Figure 3.11: AES elements profile of carburized wrought 2205
Figure 3.11 displays the AES line scan from the carburized wrought 2205 sample taken along the path shown in figure 3.10. The higher Cr and lower Ni content of the ferrite phase enables it to be distinguished from the austenite phase. In the ferrite grains, the carbon concentration is abnormally high, which indicates that a transformation to a carbide phase may have occurred in ferrite grains during carburization.
The hardness profile measured by the Swagelok Company of the same sample is shown in Figure 3.12. The hardness of material has increased from
400 HV in the core to over 1000 HV in the outer 7 μm of the sample.
27
1100
1000
900
800
700
600
500
400 Hardness (HV) 300
200
100
0 0 10 20 30 40 50 Distance from edge (um)
Figure 3.12: Carburized 2205 surface hardness
To test the assumption that the ferrite phase changed to a carbide phase during carburization, electron backscattered diffraction (EBSD) analysis of the carburized 2205 was performed by Dr. Shaghi-Moshtaghin. His results are shown in Figure 3.13. From the SEM image there are two distinct phases in the carburized 2205. The pattern obtained from one phase fits very well with a
FCC structure, which demonstrates that that is the austenite phase. On the other hand, the pattern taken from the other phase, which was the BCC phase before processing, does not match that of a BCC structure. This indicates that the carburization process has changed what was previously the ferrite phase.
28
a Figure 3.13: Plan view SEM image of carburized 2205 (a), EBSD pattern from FCC phase (b) and EBSD pattern from previous BCC phase (c).
b c X-Ray Diffraction (XRD) was applied to both non-treated and carburized 2205 by Dr. Kahn. His results are shown in Figure 3.14. The magnitude of the austenite phase’s peak shift indicates that the carburization process expanded the lattice parameter of the austenite in 2205 by 3.2%. The result of a XRD surface residual stress measurement on carburized 2205 is shown in Figure 3.15. Again, it was obtained by Dr. Kahn. Elastic constants of austenite in duplex stainless steel used in this calculation are from N. Jia et al. [23]. The magnitude of the compressive stress in the austenite phase is – 3.9±0.5 GPa, calculated by the method mentioned in appendix I.
29
(200) (110) (220) austenite
ferrite (111) (311)
(200) (331) (220) (310) (420) (222)
Figure 3.14: XRD analysis result of non-treated and carburized 2205
a b Figure 3.15: Carburized 2205 XRD residual stress measurement result, austenite phase (a), the other phase (b).
30
Although the phase the ferrite transformed into has not been identified, the
XRD analysis indicates that the phase is under compressive residual stress.
The slope of the line in Figure 3.15 indicates that the residual stress must be
compressive. Compressive residual stress would be expected because the
molar volume of eventually all iron carbide phases is greater than the molar
volume of ferrite. These are shown in Table 3.2; the volume expansion is
calculated based on the molar volume of ferrite, 7.09 cm3/mole. Using the
bulk modulus of the three carbides, the elastic stress due to the constrained
lattice expansion could be estimated.
Table 3.2: Estimated volume expansion, induced stress according to corresponding carbides Relevant Estimated Iron or Volume for unit Volume Bulk Modulus Stress Carbide Fe Expansion (GPa) (GPa) (cm3/mole)
Fe 7.09
a Fe3C 7.77 9.64% 174 5.4
b c Fe5C2 7.79 9.82% 209 6-7
d e Fe7C3 8.01 13.00% 262 11
a: average value from the experimental data of the Young’s modulus of polycrystalline
cementite [24]
b: calculated from the Fe5C2 carbide structure given in H. Faraoun’s paper [25]
31
c: the bulk modulus value from H. Faraoun’s paper [25] d: calculated from the density value from Bouchard [26]
e: the bulk modulus value from K. Henriksson et al. [27]
This can be checked with the XRD residual stress measurement result
(shown in Figure 3.15) for Fe3C, using the elastic constants from C. Jiang et al. [28]. The magnitude of the compressive stress in cementite would be – 4.5
GPa, calculated by the method mentioned in appendix I. Therefore, the magnitude of estimated stresses from the lattice expansion should be the upper limit for the residual stress, because plastic flow might occur in the carbide and reduce the value of the residual stress.
32
3.2 Tensile Testing
The samples were all machined according to ASTM specification for tension
test specimens. And samples are longitudinally mechanically polished along
the loading axis. Figure 3.16 shows the geometry of samples which have
been used for both tensile and fatigue tests (unit in inch).
Figure 3.16: Geometry of samples used for tensile and fatigue tests
Then tensile testing was conducted using MTS 810 materials testing machine
as shown in figure 3.17. The crosshead displacement rate was set as 0.25
mm/min and the data were acquired every 2 second for the entire tension test.
A gage two inches long was put on the sample to measure the strain until it
reaches 1%. After tensile data were obtained, stress vs strain curves were
ploted and the 0.2% yield stress and ultimate tensile stress were calculated.
The final elongation of sample was determined by the distance between the
two ends of fracture sample parts and measured with a vernier caliper. The
final area of fracture surface was also measured with a vernier caliper.
33
Figure 3.17: MTS 810 materials testing system machine
3.3 Fatigue Testing
The fatigue tests were also done using the same geometry for samples as shown in Figure 3.16 and employing on a MTS 810 machine. All the samples were tested under fully reversed axial tension-compression loading (R=-1).
The loading form was sinusoidal and the frequency was set at 1 Hz, 10Hz, and 20Hz as necessary depend on applied load in order to prevent heating of the specimens.
34
3.4 Charpy Testing
Charpy tests were conducted according to ASTM standard by Mr. Tuma.
3.5 Hardness Testing
Buehler Mircormet® micro hardness meter was used to conduct Vickers
hardness tests. Before testing, all the samples were polished with 0.05 μm
aluminum oxide suspension. The hardness data were attained using a 50
gram force.
3.6 Fractography
The fatigue fracture surfaces of the samples were observed using a Philips XL
30 Environmental Scanning Electron Microscope (ESEM). All samples observed by ESEM were ultrasonically cleaned in acetone for three minutes and dried by pressurized air to remove any particles and other contamination on the surface prior to loading into the chamber of the ESEM. The accelerating voltage used was set at 15 kV.
35
4. Results and Discussions
4.1 IN718 Ni-base Alloy
Tensile Testing Results
The alloy IN718 is commonly solution treated and aged to make it stronger
through precipitation. Therefore, the carburization process was applied to the
alloy IN718 samples in the solution treated and aged condition. To explore the effect of ageing, samples solution treated, but not aged were also tested. Four
conditions were subjected to uniaxial tensile testing: 1.solution treated, aged
and carburized, denoted as C, 2. solution treated, aged, denoted as NT, 3.
solution-treated, denoted as ST, 4. solution treated, aged and given the same
thermal cycle as carburization, but in an inert gas atmosphere, denoted as HT.
Table 4.1 shows the tensile test results. The fracture stress is calculated using
the fracture load divided by the final fracture area of the samples. The 0.2%
yield stress and UTS results are very resemble to those found in Y. Ono’s
study of IN718. [16]
Table 4.1: Uniaxial tensile tests result of carburized, non-treated, solution-
36
treated and heat-treated IN718 superalloy Sample 0.2% U.T.S. %Elongation %Reduction Fracture
Yield (MPa) of Area Stress
Stress (MPa)
(MPa)
C 1100 1350 21 39 2530
NT 1100 1350 24 51 2300
ST 425 870 52 61 -
HT 1155 1380 24 52 2520
1400
1200
1000 IN718 NT IN718 C IN718 HT 800 IN718 ST
600 Stress (MPa)
400
200
0 0.000 0.002 0.004 0.006 0.008 0.010 0.012 0.014 0.016 0.018 Strain
Figure 4.1: Stress-strain curves of carburized, non-treated, solution-treated
and heat-treated IN718 samples in the low strain range.
From tensile tests, it is found that the 0.2% yield stresses of carburized, non-
treated and heat-treated samples are virtually identical, all around 1100 MPa.
37
The solution-treated samples have a much lower yield stress than all aged
ones, because precipitates generated in the ageing can greatly increase the
strength of alloy by inhibiting dislocation movement.
Figure 4.2 shows the stress-strain curves of carburized, non-treated, heat-
treated and solution treated samples. It should be noted that essentially no
embrittlement after carburization was observed. Carburized specimen shows
the same ductility as the non-treated specimen. The carburized case only
constitutes about 1.5% of the sample’s volume. This observation is similar
what was found for 316L stainless steel [15]. Figure 4.3 shows the true
stress-strain curves of IN718 samples in four categories respectively.
IN718 HT IN718 C 1400 IN718 NT IN718 ST 1200
1000
800
600 Stress (MPa)
400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45 0.50 0.55 0.60 Strain
Figure 4.2: Engineering stress-strain curves of carburized, non-treated, solution
treated and heat-treated IN718 in the final fracture region. Strain obtained based up crosshead displacement.
38
True s-s curve IN718 HT True s-s curve IN718 C 1800 True s-s curve IN718 NT True s-s curve IN718 ST 1600
1400
1200
1000
800 Stress (MPa) 600
400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45 0.50 Strain
Figure 4.3: True stress-strain curves of carburized, non-treated, solution treated and heat-treated IN718 in the final fracture region. Strain obtained based up crosshead displacement.
Hardness Testing Results
Indentations are taken in the center of sample’s grip, within the area corresponding to the gage area, i.e. in the central 5 mm diameter circulous region. All samples are prepared from the grip to avoid a cyclic stress hardening or softening effect. Indentations were conducted along two arbitrarily chosen perpendicular directions, with a 25 μm distance between each other.
In both carburized and non-treated samples, several soft areas were found in the center. Figure 4.4 shows the typical soft spot phenomenon found on a
39 non-treated sample tested under a 550 MPa cyclic stress. It is noteworthy that this sample is the only non-treated sample that showed internal fatigue crack initiation, as can be seen in Figure 4.5.
one direction perpendicular direction 600
550
500
450
400
350 Hardness (HV)
300
250
200 -3000 -2000 -1000 0 1000 2000 3000 Distance (um)
Figure 4.4: Hardness data of IN718 sample’s grip
Two extremely soft spots found in the center indicate there is an inhomogeneous distribution of hardness in the IN718. The exact explanation for this phenomenon is not understood yet. However, it may be due to a casting defect.
40
a
b
Figure 4.5: Fractograph of non-treated IN718 sample, showing internal fatigue crack initiation (550 MPa)
41
Charpy Test Results
Table 4.2 lists the charpy test results of both non-treated and carburized
IN718 samples. It indicates that the impact energy of IN718 is not affected by carburization.
Table 4.2: Charpy test results of non-treated and carburized IN718 Impact Energy (N-m)
Non-treated IN718 35
Carburized IN718 33
42
Fatigue Testing Results
The fatigue tests conducted on NT and C specimens are plotted in Figure
4.6a to generate S–N curves (maximum stress in the fatigue cycle plotted vs. the number of cycles to failure). One C specimen was pulled in tension to 1% total strain before fatigue testing. The stress-strain behavior of the pre- strained sample is shown in Figure 4.6b. About 0.5% plastic strain was induced.
The fatigue test results of NT, C and pre-yielded C samples, shown in Figure
4.6a, revealed a complex fatigue behavior. At relatively high stress level (>
550 MPa, 50% of Y.S.), NT and C specimens have very similar fatigue lifetimes. At lower stress (< 550 MPa), NT specimens have longer fatigue lifetimes, but the pre-yielded C sample shows the same fatigue behavior as the NT samples. Both non-treated and pre-yielded C samples have fatigue limit at 445 MPa. In Y. Ono’s study, the IN718 fatigue limit corresponded to stress amplitude, σa, of 340 MPa when tested under an applied stress ratio (R
ratio) of 0.01. [16] According to the Goodman relation,
σ m σσaR==−1(1 − ) ……………………………………………………………..(1) σUTS 1+ R From Y. Ono’s data, σσma= , so σm of 347 MPa, σUTS=1369 MPa, 1− R
σa=340 MPa. Therefore, if their samples were tested in R=-1 case, the fatigue
limit would be 455 MPa, which match well the result of this study.
43
800 IN718 Non-treated IN718 Carburized 750 IN718 Carburized yielded
700
650
600
550 * Stress (MPa)
500
450
400
350 100000 1000000 1E7 Number of Cycles * indicate the only NT sample has internal crack nucleation
a
IN718 Carburized
1200
1000
800
600 Stress (MPa)
400
200
0 0.000 0.002 0.004 0.006 0.008 0.010 0.012 Strain
b
Figure 4.6: (a) Fatigue results for C and NT specimens. (b) Stress-strain behavior before fatigue testing for the yielded C specimen shown in (a). Fractography
44
a
Figure 4.7: SEM images of fracture surface of carburized sample (a), (b); non-treated sample (c), (d). Both the samples failed at 550 MPa stress, 50% of Y.S. (e) is the shear shape at the edge of fractured fatigue carburized sample
45
b
c
46
d
e
47
Figure 4.7 shows the typical fatigue fracture surface features of carburized
and non-treated samples. All the carburized samples have internal fatigue
crack nucleation, or more accurately, cracks initiate at the center of sample as
shown in (a) and (b). During the cyclic loading, cracks propagated towards the
surface. As the crack grows, eventually, the remaining materials can not
support the load elastically, and flow occurred, as demonstrated by the ˚45
shear structure around surface, shown in (e). In contrast, almost all the
fractured non-treated samples showed surface crack initiation, as shown in (c)
and (d). Only one sample, tested at 550 MPa, showed interior crack initiation.
Cracks initiated at the non-treated specimens’ surface, then propagated to the
center under repeated stress cycles, finally leading to an ultimate ductile
fracture. This phenomenon can be explained because of the generation of a
high compressive residual stress (about 1.9 GPa) at the surface during
carburization, which suppresses crack initiation at the surface. It should be
mentioned that both samples were tested under the same conditions: R=-1,
cyclic frequency of 20 Hz and maximum stress at 550 MPa. The fatigue life of
the carburized sample was 1.07×106 cycles, while non-treated sample fractured at 4.5×105 cycles.
From the fatigue tests result, it is clear that at the higher stresses (above 550
MPa, 50% of yield stress), carburized samples exhibited an increment in
fatigue life by a factor of about 2, compared with non-treated samples.
However, at lower stresses (under 50% of yield stress), instead of enhancing
48
the fatigue properties of IN718, carburization decreased the endurance limit.
Non-treated sample failed after 3.68×106 cycles at 500 MPa ≈45%( of yield
stress) while carburized fractured at 1.65×106 cycles. Moreover, at a stress
equal to 450 MPa, non-treated samples ran out at 10 million cycles, while
carburized sample failed at 3.5×106 cycles.
The decrease in the fatigue resistance of IN718 can be explained by a combination of casting defects and the effect of residual tensile stress in the center of samples. Usually, the carburization process causes formation of high
compressive residual stresses in the carburized case, effectively negating the
importance of stress concentrators such as grooves, scratches and other
surface defects, which cause crack initiation at the surface. However, to
counteract the high residual compression at the surface, there is a low
residual tensile stress generated in the core. The amount of this tensile stress
can be calculated from following equation:
σcompressive * Acase + σtensile * Acore= 0……………………………………………(2)
According to the available information such as case depth 15 μm, the total
gage diameter 5.08 mm (0.2 inch), and the residual compressive stress -1.9
GPa, the estimated internal tensile stress is 23 MPa. This indicates the mean stress on the interior of the carburized samples is 23 MPa.
Because of the existence of residual stresses in carburized samples, the fatigue tests are not conducted in fully reversed tension and compression mode, despite the nominal applied stress set at R=-1. According to the
49
Goodman relation,
σ m σσ aR = =− 1 (1 − ) ……………………………………………...... (1) σUTS and available data σm=23 MPa, σR=-1=445 MPa, σUTS=1.3 GPa, we can calculate that the stress amplitude is 437 MPa when a 23 MPa mean stress exists. Thus, the effect of the 23 MPa mean stress is to reduce the stress amplitude at the endurance limit by 8 MPa.
In addition, typical fatigue elements like intrusion, extrusion (indicated by an
arrow in Figure 4.8) and persistent slip bands (PSB) were found at the fatigue crack initiation sites of IN718 fatigue testing bars. Figure 4.8 was obtained from a HT sample fatigue fractured at 450 MPa. Figure 4.8 (a) and (b) show
the lateral view image of the surface crack origin under different
magnifications. Figure 4.8 (c) is the plan view at the fatigue origin of the same sample. Figure 4.8 illustrates the formation of fatigue cracks in the IN718 alloy. The PSBs pile up at a certain position due to the cyclic reversed stress,
and at the surface, intrusions and extrusions, which are indicated by arrows in the images, are generated, which make the stress concentrations at the specimen surface even greater. Finally, the local stress exceeds a critical value and cracks initiate.
50
a
Figure: 4.8: SEM images of a fatigue fractured HT IN718 sample, (a) and (b) lateral view, (c) plan view
b
51
c For the fully-reversed fatigue tests performed in this study, the fatigue lifetimes will be a function of the stress amplitude Δσ/2 as given by equation (3). [2]
∆σ ' b = σ ff(2N ) 2 …………………………………………………………………….(3)
σf’ is the fatigue strength coefficient, and b is known as the fatigue strength or
Basquin exponent. The fatigue lifetimes of carburized and non-treated IN718
samples are plotted using a log-log scale in Figure 4.9. The Basquin
exponent, b, is a strong indicator of the magnitude of the increase in lifetime
with decreased fatigue stresses. The lower b leads to greater improvements in
fatigue life as the stress decreases. The fatigue strength coefficient σf’ is
typically a good approximation for the actual fracture stress. Table 4.3 gives
the values of b and σf’ for carburized and non-treated IN718.
52
5000 IN718 Non-treated IN718 Carburized Linear fit of IN718 NT Linear fit of IN718 C Stress (MPa)
400 1 10 100 1000 10000 100000 1000000 1E7 Number of Cycles
Figure 4.9: Log-log plot of carburized and non-treated IN718 fatigue results
Table 4.3: Basquin parameters for carburized and non-treated IN718 samples Sample Fatigue Strength Fatigue Strength
exponent, b coefficient, σf’
IN718 -0.15 4500 MPa
Carburized
IN718 -0.11 2600 MPa
Nontreated
A summary of the results regarding fatigue testing of the IN718 alloy are as follows:
• There are some soft regions in the central (gage) portions of test
sample for both non-treated and carburized samples, due to the
53
presence of possible casting defects.
• The carburization process does not change the microstructure and
mechanical properties in the core. However, the fatigue strength
coefficient increases after carburization, so does fracture stress.
• Because the hard and highly compressed case suppresses surface
crack nucleation, fatigue cracks in carburized samples are generated in
the center. On the other hand, non-treated samples generally exhibit
surface crack initiation, unless the interior strength is extremely low.
• Due to the effect of residual tensile stress generated in the core to
balance compressive stress in the case, carburized samples have a
fatigue limit that is about 8 MPa lower than non-treated samples when
interior defects existed. Presumably, if these defects were eliminated,
the carburization would improve the fatigue life due to the effect of the
carburized case upon fatigue strength.
54
4.2 A286 Fe-base Alloy
Tensile Testing Results
The A286 superalloy samples in both carburized and non-treated conditions
were tested under uniaxial tension. There were three batches of samples including both a carburized and a non-treated specimens in each batch. Batch
A samples had rougher circumferential grooves perpendicular to the specimen axis. Batch B samples’ surface grooves were finer, while still circumferential.
Batch C samples had no circumferential grooves but longitudinal scratches along the gage. The tensile test results are listed in the Table 4.4. The fracture
stress is calculated using the fracture load divided by the final fracture area of
the samples.
Table 4.4: Uniaxial tensile tests data of carburized, non-treated A286
superalloy
Batch Sample 0.2% U.T.S. %Elongation %Reduction Fracture
Yield (MPa) of Area Stress
Stress (MPa)
(MPa)
A Carburized 855 1135 14 38 - Non-treated 880 1165 22 35 -
B Carburized 855 1130 22 34 - Non-treated 820 1135 23 50 -
C Carburized 785 1110 16 39 1460 Non-treated 805 1115 18 49 1650
55
A286 NT 1000 A286 C
800
600
400 Stress (MPa)
200
0 0.000 0.002 0.004 0.006 0.008 0.010 Strain
Figure 4.10: Stress-strain curves of batch C carburized and non-treated A286 samples in the low strain range
Figure 4.10 shows the 0.2% yield stresses of carburized and non-treated
A286 were approximately the same.
56
A286 C new A286 NT new
1200
1000
800
600 Stress (MPa) 400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 0.30 Strain
Figure 4.11: Stress-strain curves of batch C carburized, non-treated A286 samples
True S-S A286C True S-S A286NT 1400
1200
1000
800
600 Ture stress (MPa) 400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 Ture strain
Figure 4.12: True stress-strain curves of carburized, non-treated A286 samples
57
From both Table 4.4 and Figure 4.11, it is observed that all of the A286 samples, independently of surface finish, had the same value of ultimate tensile stress, about 1.1 GPa. Consistent with the pervious observations, carburized specimens show slightly lower fracture strain than non-treated samples, which indicates little or no embrittlement due to the carburization.
Figure 4.12 shows the true stress-strain curves of carburized and non-treated
A286 samples.
The fracture surfaces of batch C samples, both non-treated and carburized tensile tested were examined with an ESEM. The Figure 4.13 shows the features on a lateral surface near the fracture surface. The non-treated sample showed more extensive necking behavior and a number of flow lines near the neck. There were numerous micro-cracks instead of flow lines on the carburized sample, where the surface appears to have become less ductile after the carburization process. These images indicate that non-treated samples have slightly higher surface ductility and also fractured in a slightly more ductile way.
58
a
Figure 4.13: Side view image of A286 rod samples fractured in tension, (a) and (b) non-treated; (c) and (d) carburized.
59
b
c
60
d
Charpy Test Results
Table 4.5 lists the charpy test results for both non-treated and carburized A286 samples. It indicates that the fracture toughness of A286 is not affected by carburization.
Table 4.5: Charpy test results of non-treated and carburized A286 Impact Energy (N-m)
Non-treated A286 53
Carburized A286 51
61
Fatigue Testing Results
Fatigue properties of carburized and non-treated A286 samples in batch C were investigated. The fatigue testing results are shown in Figure 4.14. The fatigue testing results illustrates that:
1. Carburized samples have approximately the same fatigue life as the non- treated ones.
2. Both carburized and non-treated samples show the same fatigue endurance limit of 350 MPa.
3. As shown below the carburization process affected the fatigue strength coefficient and exponent of the A286 superalloy.
The Basquin analysis was applied to the fatigue results for carburized and non-treated A286, as shown in Figure 4.15. The values of b and σf’ attained are listed in Table 4.6.
Table 4.6: Basquin parameters for carburized and non-treated A286 samples Sample Fatigue Strength Fatigue Strength
exponent, b coefficient, σf’
A286 -0.098 1700 MPa
Carburized
A286 -0.076 1200 MPa
Nontreated
62
A286 Carburized Batch C 600 A286 Non-treated Batch C
550
500
450 Stress (MPa)
400
350
10000 100000 1000000 1E7 Number of cycles
Figure 4.14: The batch C carburized and non-treated A286 fatigue life.
A286 Carburized Batch C 2000 A286 Non-treated Batch C Linear fit of A286 C Linear fit of A286 NT 1500
1000 Stress (MPa)
500
1 10 100 1000 10000 100000 1000000 1E7 Number of Cycles
Figure 4.15: Log-log plot of carburized and non-treated A286 fatigue results
63
Fractography
a
Figure 4.16: SEM images of fracture surface of non-treated fatigue samples fractured at 380 MPa (a), (b); 510 MPa (c), (d), (e).
64
b
c
65
d
e
66
a
Figure 4.17: SEM images of the fracture surfaces of carburized sample fatigue fractured at 380 MPa (a), (b); 470 MPa (c), (d).
67
b
c
68
d
Figure 4.16 shows the typical non-treated A286 fracture surfaces. All the non-
treated fatigue bars tested under different stress levels exhibited surface
crack initiation. However, the number of crack originating at the surface
depended upon the magnitude of the stress. At lower stress, 380 MPa (45% of yield stress), there was only one crack initiation site on the edge. All the cracks grew out from that single origin. For the sample tested under 510 MPa
(63% of yield stress), it is obvious that at least two crack initiation sites were
active, located at the upper left side and upper right side (shown by the arrow)
respectively in images (c), (d) and (e). Multiple crack origins occur because at
higher applied stress, it is easier to generate cracks at numerous surface
defects such as inclusions, particles and grooves on the surface.
69
On the other hand, the carburized samples behaved differently. The general
observation (Figure 4.17) is that the number of crack origins was not affected
by the stress level. There was always only one crack initiation site for the
carburized samples tested at 380 MPa (45% of yield stress) and 470 MPa
(60% of yield stress). However, similar to the non-treated samples, the
carburized samples had cracks initiate at their surfaces. One possible explanation is that, after carburization process, the surface of sample became less ductile, which is observed in the tensile testing. Therefore, the harder but less ductile surface became favorable for the crack initiation during cyclic loading. This indicates why the carburized samples had fatigue lives similar to
those of the non-treated samples.
XRD Residual Stress Analysis
Compared with the 1.9 GPa residual stress of carburized IN718 Ni base alloy
in this study, the 3.3 GPa residual compressive stress generated in carburized
A286 is higher. However, the carbon concentration at the materials surface for
both IN718 and A286 is the same, ~12 at.%. This result can be attributed to
the difference of effects of carbon content on lattice expansion for Ni and
austenite.
From the study of C. Roberts [29], the relation between carbon concentration
and austenite lattice parameters can be given by a (Å) = 3.548+0.044xc; xc is
weight percent of carbon. Therefore, the relation also can be equally
expressed by a=3.548+0.010xc’; xc’ is atomic percent of carbon.
70
From L. Zwell et al.[30], the changes in lattice parameter with variation in carbon concentration could be expressed by a (Å) = 3.528+0.0074x.
Therefore, the same amount of carbon atoms could induce lattice expansion in austenite approximately 1.4 times more than in nickel. This is the reason why carburized IN718 has lower residual stress than carburized A286.
71
4.3 2205 Duplex Stainless Steel
Tensile Testing Results
Wrought 2205 duplex stainless steel samples both in the non-treated,
carburized and subjected to the same thermal cycle as employed for
carburization, but in an inert atmosphere (indentified as heat treated)
conditions were tested under uniaxial tension. The tensile test results are
listed in Table 4.7. The yield stress and ultimate stress of the heat-treated sample are slightly higher than the carburized material and both values for the carburized sample are 18% higher than those of the non-treated sample. This indicates the heat treatment led to a strengthening similarly to previous observation from LTCSS processing. The fracture stress is calculated using the fracture load divided by the final fracture area of the samples. It is found
that the tensile properties of the non-treated samples are similar to those
reported data in Mateo’s study. [19]
Table 4.7: Uniaxial tensile tests data of non-treated, carburized and heat-treated
2205 superalloy
Sample 0.2% Yield U.T.S. %Elongation %Reduction Fracture
Stress (MPa) of Area Stress
(MPa) (MPa)
Non-treated 565 765 38 88 3030 Carburized 670 905 28 76 1910 Heat-treated 680 920 - - -
72
2205NT wrought 800 2205C wrought
700
600
500
400
Stress (MPa) 300
200
100
0 0.000 0.002 0.004 0.006 0.008 0.010
Strain
Figure 4.18: Stress-strain curve of carburized and non-treated 2205 samples
in the low strain range
Figure 4.18 shows that the carburization process markedly increased the
0.2% yield stress of 2205 samples, from 565 MPa to 670 MPa.
73
2205NT wrought 1000 2205C wrought UTS= 910 MPa 900
800
700 UTS= 770 MPa 600
500
400 Stress (MPa)
300
200
100
0 0.0 0.1 0.2 0.3 0.4
Strain
Figure 4.19: Stress-strain curves of carburized, non-treated 2205 samples
True s-s curve 2205 NT True s-s curve 2205 C 1200
1000
800
600 Stress(MPa) 400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35
Strain
Figure 4.20: True stress-strain curves of carburized, non-treated 2205 samples
74
From both Table 4.3 and Figure 4.19, it is clear that after carburization, 2205 sample became much stronger, with a 140 MPa increment in the UTS.
Consistent with the previous observations, the carburized specimens exhibits a slightly lower fracture strain than non-treated one. However, the toughness, calculated by taking the integral underneath the stress-strain curve, of the carburized 2205 is slightly increased. The true stress-strain curve is shown in
Figure 4.20.
The fracture surfaces of both non-treated and carburized tensile samples were observed with a SEM, as shown in Figure 4.21. The non-treated sample showed a larger reduction of area and a number of “flow lines” near the necking. On the other hand, there are numerous micro-cracks rather than flow lines on the carburized sample, which indicates the surface become became less ductile after carburization. These images indicate that the non-treated sample had greater surface ductility.
75
a
Figure 4.21: Tension fracture surface and neck area of 2205 samples, (a) and
(b) non-treated; (c) and (d) carburized
76
b
c
77
d
Charpy Test Results
Table 4.8 lists the charpy test results of both non-treated and carburized 2205 samples. It shows that the impact energy of 2205 decreased by a factor of 2 after carburization.
Table 4.8: Charpy test results of non-treated and carburized 2205 Impact Energy (N-m)
Non-treated 2205 320
Carburized 2205 160
78
Fatigue Testing Results
The fatigue properties of carburized and non-treated wrought 2205 samples
were investigated. Testing results are shown in Figure 4.22. The fatigue
testing results illustrates that:
1. 2205 is a very fatigue-resistant alloy, which can endure more than 107
cycles at 75% of its 0.2% yield stress in fully reversed tension-
compression testing.
2. The carburization process significantly improves alloy 2205’s fatigue
resistance.
(a) carburized samples obtain an approximately 30 MPa increment in
fatigue endurance limit compared with non-treated samples.
(b) at all stress level, carburized samples show more than a 10 times
improvement compared with non-treated ones (204,000 cycles vs 15,000
cycles at 510 MPa, and 2.5×106 cycles vs 105 cycles at 480 MPa, run out
107 vs 3×106 at 450 MPa)in fatigue life.
3. For non-treated samples, at a high stress condition (> 450 MPa), surface crack initiation is dominant as shown in Figure 4.25 a, b. At low stress levels
(< 450 MPa), cracks are nucleated at preexisting oxide inclusions in the materials, and then propagate towards the surface as shown in Figure 4. 26.
4. Oxide inclusions are seen to play a crucial role in crack initiation in all carburized samples for example a shown in 4.25 c, d. Subsurface “fish eye” crack formation was found in those fractured carburized samples.
79
2205NT wrought 600 2205C wrought
550
500
450
Stress (MPa) 400
350
10000 100000 1000000 1E7 Number of Cycles
Figure 4.22: S-N curve shows the carburized and non-treated 2205 samples fatigue life
2205 from this study 2205 Mateo's group data
520
480
440
400 Stress (MPa)
360
320
100000 1000000 1E7 Number of Cycles
Figure 4.23: Comparison S-N curves of the non-treated 2205 samples in our group and L samples in Mateo’s group [18]
80
From the Figure 4.23, it is interesting to find that 2205 samples tested in different groups have the same fatigue limit of approximate 430 MPa. At a maximum stress amplitude of 440 MPa, the result Mateo’s group matches well that of this study. However, at higher stresses, Mateo’s group samples fatigue lives are about 6-8 times higher than those of this study. This outcome can be attributed to the defects found in the samples from this study which are discussed previously.
1000
2205 Non-treated 2205 Carburized 800 Linear fit of 2205 C Linear fit of 2205 NT
600 Stress (MPa)
400 1 10 100 1000 10000 100000 1000000 1E7 Number of Cycles
Figure 4.24: Log-log plot of carburized and non-treated 2205 fatigue results
The Basquin analysis was applied to the fatigue results for carburized and non-treated 2205, as shown in Figure 4.24. The values of b and σf’ are listed in Table 4.9.
The fatigue strength coefficients that were determined are anomalous because of their very low values.
81
Table 4.9: Basquin parameters for carburized and non-treated 2205 samples Sample Fatigue Strength Fatigue Strength
exponent, b coefficient, σf’
2205 -0.034 790 MPa
Carburized
2205 -0.026 645 MPa
Nontreated
82
Fractography
a
Figure 4.25: SEM images of fracture surface of sample fatigue fractured at
480 MPa: non-treated (a), (b); carburized (c), (d); (e) shows the XEDS spectrum and chemical composition of inclusion particle.
83
b
c
84
d
O 63.8 Mg 0.7 Al 20.4 Si 0.5 Ca 2.7 Ti 0.5 Cr 3.1 Mn 0.5 Fe 6.7 Ni 0.6 Mo 0.3
e
85
a
Figure 4.26: SEM images of fracture surface of non-treated sample fatigue
fractured at 440 MPa (a), (b); (c) shows the XEDS spectrum of inclusion
particle.
86
b
c
XRD Residual Stress Analysis The magnitude of compressive residual stress in the austenite phase in carburized 2205 is 3.9±0.5 GPa, which is slightly higher than that in the carburized 2205. Though the slight different may be due to error in measurement, from the surface hardness testing results (Figure 3.7 and
87
Figure 3.12), the carburized 2205 has a harder surface than the carburized
A286. It is reasonable to assume that the possible carbide phase, which replaces the ferrite phase, constrains the flow of the carburized austenite phase. Consequently, the carburized austenite phase in 2205 can reach a higher stress than its counterpart in A286.
88
5. Conclusions
All the current research demonstrated that the Low Temperature Carbon
Supersaturation (LTCSS) process can be successfully applied to alloys IN718,
A286 and 2205 stainless steel. Tensile and fatigue testing were conducted on these alloys. It is observed that the carbon-rich case with both a large compressive residual stress and significantly increased hardness plays an important role in the fatigue behavior of these alloys.
In the carburized IN718 samples, the carburization-induced residual compressive stresses eliminate surface-initiated fatigue cracks. However, since defects (possibly from casting) exist in the center of the samples which are not protected by the compressive stresses in the carburized case, the fatigue lifetimes were not improved. In fact, at low applied stresses, the lifetimes were slightly decreased due to effect of the residual tensile stress in the core.
The fatigue life time of A286 alloy was slightly affected by the carburization process, which is not completely understood. Carburized samples attained modest improvements in fatigue life. However, carburization effectively suppressed the formation of multiple crack nucleations at the surface.
Consistent with the previous observation in 316L stainless steel, carburization
89 significantly increased the fatigue resistance of 2205 duplex stainless steel.
Carburized 2205 samples exhibited increases both in fatigue endurance limit
(improved from 420 MPa to 450 MPa) and fatigue lifetimes compared to non- treated samples (as much as 27 times) at the same stress amplitude. Surface crack initiation was suppressed after carburization. All the cracks in the carburized samples generated from oxide inclusions in the center of the samples. In addition, subsurface “fish eye” crack formation was found on several fractured surfaces.
For all three alloys, the fatigue strength exponents and fatigue strength coefficients derived from the Basquin equation, along with the 0.2% Y.S.,
UTS, and % elongation at fracture are summarized in Table 5.1. For all three alloys, carburization increased the fatigue strength coefficient and decreased the fatigue strength exponent. Carburization also decreased the percentage of elongation at tensile fracture.
90
Table 5.1: Basquin parameters for carburized and non-treated 2205, IN718 and A286 Sample Fatigue Fatigue 0.2% Y.S. U.T.S. %
Strength Strength (MPa) (MPa) Elongation
exponent coefficient
, b (MPa) , σf’ (MPa)
IN718 -0.15 4500 1110 1355 21
Carburized
IN718 -0.11 2600 1110 1355 22
Nontreated
A286 -0.098 1700 830 1125 17
Carburized
A286 -0.076 1200 835 1140 21
Nontreated
2205 -0.034 790 670 905 28
Carburized
2205 -0.026 640 565 765 38
Nontreated
91
Future Work
1. Continue the study of the high cycle fatigue behavior of Ni-base IN718 with
the aim of eliminating the effect of interior defects. Use methods such as
choosing better cast materials and designing the geometry of the test
specimen to reduce the defect density in the samples.
2. To study the fatigue properties of alloy IN718, A286 and 2205 under
corrosive environments (corrosion fatigue testing).
3. To study these alloys using other fatigue testing procedures, such as
tension-tension fatigue, rotating bending fatigue.
4. Combine the study of both the initiation and propagation stage of fatigue
cracks to understand fatigue behavior of these alloys more completely.
5. Fatigue testing of other alloys, which have been successfully carburized by
the LTCSS, such as martensitic steels 13-8, 15-5, 17-4, and the ferritic
steel E-Brite.
92
APPENDIX I
Residual Stress Estimation in Thesis
The residual stresses of all samples were determined by the standard XRD
sin2ψ technique. [A1]
EEdd−− aa σ = in= in 22 ……………………………(1) (1++νψ) sin dann(1 νψ) sin
σ, denotes the residual stress; E, Young’s modulus; ν, Poisson ratio; ψ,
various tilt angles; di, lattice spacing of a particular hkl plane at various tilt
angles; dn, lattice spacing of the same hkl plane at ψ=0. And the lattice
parameters attained from di and dn are ai and an.
According to De Wit’s study on the diffraction elastic constants of cubic
polycrystal, [A2] the ν and E parameters can be represented E G = ……………………………………………………………………...……..(2) 1+ν by the effective shear modulus of the effective medium, G, as shown in
equation 2. The value for G is equal to 1/S2, where S2 is the second
diffraction elastic constant. Therefore, the residual stress can be given by
slope 2 σ = × …………….………………………………………………..(3) intercept S 2hkl
ddin− where slope is defined by ; intercept is dn. Figure A1, shows how sin2 ψ residual stress is calculated in carburized IN718.
93
Figure A1: XRD lattice parameters for the (420) peak of carburized IN718 as
a function of sin2ψ.
According to Kroner’s original cubic equation for the effective shear stress
modulus of cubic polycrystal
GGG323−αβγ − −=0 …………………………………………………….……..(4) where α=α2-α1, β=β2-β1
And α1, α2, β1, β2, γ are given by De Wit [A2] as follows:
3 '' ' '' α1 =34 κ + µ +−Γ 3( µµ) 8{ }
1 ' '' α2 =(23 µµ + ) 5 3 '' ' '' β1 = κµ +3( µ −Γ µ) ………………………………………………….(5) 4
3 ' '' ' '' β2 =(6 κµ ++ 9 κµ 20 µ µ ) 40 3 γ= κµ' µ '' 4
94
Γ was introduced to describe the diffraction effects of a cubic polycrystal and
given by
Γ=(h2k2+k2l2+l2h2)/(h2+k2+l2),
and the cubic bulk modulus and shear moduli are
1 κ =(cc + 2) 3 11 12 1 µ ' =()cc − ……………………………………………………………………(6) 2 11 12 '' µ = c44
95
Appendix II
Effect of Carburization on the Tensile Behavior of 316L Wire
Tension tests were conducted on three 316L stainless steel wires. The fracture mechanisms associated with the wires before and after carburization were determined.
Materials and Methods
The 316L wires had diameters of 76 um, 127 um, 203 um.
The same MTS 810 system described in the thesis was used to conduct tensile test on these wires.
A Hitachi S4500 Scanning Electron Microscopy was used to observe the fracture surfaces of both non-treated and carburized 316L wires.
Results
The Figure A2 shows the tension test results of 316L wires, both non-treated and carburized. It is surprising to find that for the 76 um and 127 um wires, after carburization, the wires exhibited increases in both fracture stress and fracture strain. The 203 um wire shows a stress but lower fracture strain.
The strengthening result could be attributed to annealing effect of the heat
96
treatment during the carburization process. When thin wires with 76 um and
127 um diameters were drawn from bar material, to prevent the wires from
fracture, wires were annealed succeeding draw process several times. In
other words, wires were manufactured under cycles of drawing, then
annealing. Probably, these wires were not annealed after drawn to be as thin
as 76 or 127 um. Therefore, the thermal treatment in the carburization
process annealed these wires, and made them stronger.
3000 2000 UTS= 2.7 GPa UTS=1.8GPa 1800 2500 316C 003 1 1600 316NT 003 1 1400 316C 005 1 2000 316NT 003 2 316C 005 2 1200 316NT 005 1 UTS= 1.7 GPa 316NT 005 2 1500 1000 UTS=1.17 GPa Stress(MPa)
Stress (MPa) 800 1000 600
400 500 200
0 0 0.00 0.01 0.02 0.03 0.04 0.05 0.06 0.00 0.01 0.02 0.03 0.04 0.05 0 Strain Strain a b
Figure A2: The engineering stress-strain curves of non-treated and carburized
316L wires. a) 76 um thick wires; b) 127 um thick wires; c) 203 um.
97
316C 008 1 316C 008 2 UTS=0.96 GPa 1000 316NT 008 1 316NT 008 2
800
600 UTS=720MPa Stress (MPa) 400
200
0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Strain
c
The fractography of the wires in Figure A3 shows that the non-treated wires all fractured in a very typical ductile way, a) and b). Under tension, necking occurs, and then stresses concentrate at the necking area, generating more severe necking. Finally a shear failure occurred and the fractured surface showed numerical dimples.
The carburized wires show a complex fracture behavior. For both 76 um and
203 um thick wires, a combination of fracture mechanisms is found. Near the edge of samples, where the carburized case is, the surfaces show a faceted structure. No dimples are found. However, along radius from edge to the center, more and deeper dimples occur in the core, which indicates a more ductile fracture mechanism.
98
a
Figure A3: Fracture surface of non-treated 316L wires a) 76 um, b) 203 um; and carburized wires c) 76 um, d) 203 um.
b
99
c
d
100
High Cycle Fatigue of Carburized 316L with Different Surface Finishes
Four batches of 316L fatigue bars with different circumferential groove spacings were prepared. The distances between grooves on the surface were
3.2 mil, 6.4 mil, 12.5 mil and 25 mil, named as #32, #64, #125 and #250, respectively. After the surface finishes were prepared, those samples were carburized according to the same technology discussed in the thesis. Tension and fatigue tests were conducted on these samples.
The tensile test results are listed in the following Table A1.
Table A1: Uniaxial tensile tests conducted on carburized 316L stainless steel samples having different surface finishes Sample 0.2% Yield U.T.S. %Elongation %Reduction
Stress (MPa) of Area
(MPa)
#32 316L 560 740 33 65
#64 316L 560 740 35 70
#125 316L 560 740 38 75
#250 316L 560 740 35 70
101
700 Yield Stress=560MPa 600
316C 32# 500 316C 63# 316C 125# 400 316C 250#
300 Stress(MPa)
200
100
0 0.000 0.002 0.004 0.006 0.008 0.010 Strain
Figure A4: Stress-strain curves of different surface finishes carburized 316L
samples in the low strain range.
From the Figure A4, it can be concluded that different surface finishes have
no effect on the yield stress in tension. Four stress-strain curves replicate each other in low strain region. All the specimens have yield stress of 560
MPa.
102
316C 32# 316C 63# 800 UTS=740MPa 316C 125# 316C 250# 700
600
500
400 Stress(MPa) 300
200
100
0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 Strain
Figure A5: Stress-strain curves of different surface finishes carburized 316L samples.
From the complete stress-strain curves shown in Figure A5, the similar ultimate tensile stress (UTS) and 0.2% yield stress values are observed for each of the four surface finishes. However, a difference in the fracture strain was observed. No relationship between the fracture strain and the surface finish was found. The sequence from smallest to largest fracture strain was
32#, 250#, 63# and 125#.
Fatigue Testing Results
103
Fatigue properties of carburized 316L samples with different surface finishes
were investigated. Fatigue test results are showed in Figure A6.
Figure A6 shows that samples with all four surface finishes from ran out at
400 MPa stress amplitude (70% of yield stress). The 32# and 250# samples
ran out at 425 MPa (75% of yield stress). That is a very high fraction of the
yield stress.
The difference in fatigue life between samples with different surface finishes is
small. Therefore, that conclusion can be drawn that fatigue properties of
carburized 316L stainless steel are not affected by the differences in surface
finish evaluated in this test program.
316-#63 500 316-#250 316-#125 316-#32
450 *
400 Stress(MPa)
350
300 1000 10000 100000 1000000 1E7 Number of cycles
Figure A6: S-N curve shows the different surface finishes carburized 316L fatigue life.Tests conducted with a load ratio, R=-1.
104
MP-98T Cobalt Base Superalloy Fatigue Testing
SPS MP98T superalloy, which was developed by SPS Technologies, exhibits a combination of strength, toughness, and corrosion resistance not previously available in a fastener material. These fasteners offer users a minimum tensile strength of 1240 MPa. Typical fracture toughness values for MP98T specimens tested per ASTM E 1820 exceed 220 MPa-m1/2. The strength level and fracture toughness make them well suited for critical aerospace engine and airframe applications. [A3]
The chemical composition of the MP98T alloy in this study is in Table A2.
Table A2: Chemical Composition of MP98T Composition Wt.%
Co 36
Ni 25
Cr 19
Fe 9
Mo 7
Ti 3
Nb 0.6
Al 0.2
MP98T samples with the same geometry as described in section 3.2 were prepared. Tension and full-reversed tension-compression fatigue tests were conducted on both non-treated and carburized samples.
105
The tensile test results are listed in Table A3.
Table A3: Uniaxial tensile tests conducted with the MP98T superalloy in non- treated and carburized conditions Sample 0.2% Yield U.T.S. %Elongation %Reduction
Stress (MPa) of Area
(MPa)
Non-treated 1200 1310 - -
Carburized 1330 1440 19 34
The tensile test results show that the carburization process gave the MP98T a
10% increment in its yield stress, from 1200 MPa to 1330 MPa, as shown in
Figure A7.
MP98T NT 1600 MP98T C
1400
1200
1000
800
Stress (MPa) 600
400
200
0 0.000 0.002 0.004 0.006 0.008 0.010 Strain
Figure A7: Stress-strain curves of non-treated and carburized MP98T samples in the low strain range.
106
Figure A8 shows the complete stress-strain curve of non-treated and
carburized MP98T. Compared to the non-treated sample, the carburized exhibits a 130 MPa higher UTS. However, the final fractured strain decreased from 0.25 to 0.17 after carburization.
Because of the lack of samples, only six fatigue tests were able to be performed. The results of the fatigue testing are shown in Figure A9.
MP98T C 1600 MP98T NT
1400
1200
1000
800
600 Stress (MPa)
400
200
0 0.00 0.05 0.10 0.15 0.20 0.25 Strain
Figure A8: Stress-strain curves of non-treated and carburized MP98T
samples.
107
800 MP98T NT MP98T C 700
600
500
400 Stress (MPa) 300
200
100 10000 100000 1000000 1E7 Number of Cycles
Figure A9: The fatigue testing results for non-treated and carburized MP98T.
Tests conducted with a load ratio, R=-1.
108
303 austenitic stainless steel fatigue testing
Alloy 303 is a non-magnetic, austenitic, free-machining stainless steel that is a modification of the basic 18% chromium -9% nickel stainless steel specially designed to exhibit improved machinability while maintaining good mechanical and corrosion-resistance properties.
303 stainless steel samples with the same geometry as described in section
3.2 are prepared. Tension and full-reversed tension-compression fatigue tests were conducted on both non-treated and carburized samples.
The tensile test results are listed in Table A4.
Table A4: Uniaxial tensile test results for non-treated and carburized 303 stainless steels Sample 0.2% Yield U.T.S. %Elongation %Reduction
Stress (MPa) of Area
(MPa)
Non-treated 500 805 40 63
Carburized 560 800 41 56
The tensile test results show that the carburized 303 sample exhibited a 10% increment in its yield stress from 500 MPa to 560 MPa, compared with a non- treated sample. The stress-strain curves for non-treated and carburized 303
109 stainless steel in the low strain range are shown in Figure A10.
The complete stress-strain curves contained in Figure A11, show that the carburized 303 sample has the same UTS as the non-treated sample, but a lower fracture strain.
303 NT 700 303 C
600
500
400
300 Stress (MPa)
200
100
0 0.000 0.002 0.004 0.006 0.008 0.010 Strain
Figure A10: Stress-strain curves for non-treated and carburized 303 samples in low strain range.
110
303 NT 303 C
800
600
400 Stress (MPa)
200
0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Strain
Figure A11: Stress-strain curves for non-treated and carburized 303 samples
Figure A12 shows that the carburization process improves the fatigue limit of
303 stainless steel from 300 MPa in its non-treated condition to 390 MPa when it is carburized. At 450 MPa, the fatigue life is enhanced by a factor of
120, from 6,000 cycles to failure for non-treated to 740,000 cycles to failure for carburized 303 stainless steel.
111
303 NT 303 C 500
450
400
350
Stress (MPa) 300
250
200 1000 10000 100000 1000000 1E7 Strain
Figure A12: The fatigue testing results for non-treated and carburized 303 stainless steel samples. Tests conducted with a load ratio, R=-1.
112
AL6XN super-austenitic stainless steel fatigue testing
Among a variety of stainless steels, the AL6XN alloy is distinguished by the extreme stability of its austenite, which does not transform to martensite over a wide range of temperatures and amounts of plastic deformation. While
AL6XN possesses an excellent formability typical for traditional metastable austenitic steels, there is no plasticity-induced martensitic transformation due to cold working. Therefore, the alloy is commonly referred to as a superaustenitic stainless steel. The high levels of chromium, nickel, molybdenum, and nitrogen provide an excellent resistance to chloride corrosion and stress corrosion cracking. [A4]
Table A5: Chemical Composition of AL6XN Composition Wt.%
C 0.016
Ni 23.93
Cr 20.34
Fe Balance
Mo 6.25
Si 0.29
N 0.226
Cu 0.25
AL6XN stainless steel samples with the same geometry as described in
113
section 3.2 were prepared. Tension and full-reversed tension-compression
fatigue tests were conducted on both non-treated and carburized samples.
The tensile test results are listed in Table A6.
Table A6: Uniaxial tensile test results for non-treated and carburized AL6XN stainless steels Sample 0.2% Yield U.T.S. %Elongation %Reduction
Stress (MPa) of Area
(MPa)
Non-treated 355 780 41 86
Carburized 335 780 40 81
Different from other materials studied, the tensile test results show that the
non-treated AL6XN has a slightly higher 0.2% yield stress than the carburized
sample, 355 MPa vs 335 MPa, as shown in Figure A13. Note that initial plastic
yielding of the non-treated sample begins at a lower stress than the
carburized sample.
114
500 AL6XN NT AL6XN C
400
300
Stress (MPa) 200
100
0 0.000 0.002 0.004 0.006 0.008 0.010
Strain
Figure A13: Stress-strain curve of non-treated and carburized AL6XN
samples in the low strain range
The complete stress-strain curves in Figure A14, show that the carburized
AL6XN sample has the same UTS as the non-treated sample and both have similar fracture strains.
115
AL6XN NT AL6XN C
800
600
400 Stress (MPa)
200
0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Strain
Figure A14: Stress-strain curves of non-treated and carburized AL6XN
samples
The yield stress of AL6XN is only a small fraction, i.e. 45%, of its UTS, (350
MPa vs 780 MPa). Figure A15 shows that the non-treated AL6XN has a fatigue limit of 290 MPa, which is close to it yield stress. One non-treated
Al6XN sample failed after only 30,000 cycles at 315 MPa. After carburization, the fatigue limit of the AL6XN increased to 315 MPa, although one sample of carburized AL6XN failed after 4 million cycles at 297 MPa.
116
AL6XN NT AL6XN C 320
300
280
260
Stress (MPa) 240
220
200 10000 100000 1000000 1E7 Strain
Figure A15: The fatigue testing results obtained from non-treated and carburized AL6XN samples.
117
A study of the effect of residual stress on the
hardness of carburized A286
The aim of this study
After LTCSS processing, pronounced surface compressive residual stress is induced by the expansion of the metal’s lattice parameters. This study tries to understand the effect of surface compressive residual stress on the surface hardness of alloy A286. The aim of the study is to compare material’s surface hardness before and after residual stress release. Residual stress was eliminated by yielding the samples in tension.
Materials and methods
One time carburized flat A286 tension bars were provided by the Swagelok Company. The shape of sample is shown in Figure A16.
Figure A16: Photograph of carburized A286 flat bar
The Scintag X-1 advanced X-ray diffractmeter was used to determine the surface residual stress. Tension testing was conducted on a computer-
118
controlled INSTRON Model 1361. Buehler Mircormet® was used to measure the material’s Vickers hardness using 25 gram force. Before testing, 0.5-1μm was removed from all of the surfaces by slight polishing with P4000 grit papers. All the diagonal lengths of indents were measured employing a FEI Nova nanolab 200 dual beam focused ion beam system.
Experimental Results
1. Tension testing Uniaxial tension test data was obtained by the INSTRON Model 1361 machine. The result is shown in Figure A17. Because of lack of a suitable gauge, only displacement of the machine’s crosshead was recorded.
400
350
300
250
200 Stress (MPa) 150
100
50
0 0.0 0.5 1.0 1.5 2.0 2.5 Displacement (mm)
Figure A17: Yield curve of carburized A286 flat tension bar
2. Surface residual stress analysis The magnitude of the surface stress was determined by measuring the peak positions after tilts of 0, 10, 20 and 30 degree. Figure A18 represents the
119 results of surface residual stress measurement before and after yielding. It indicates that the carburization process induces a compressive stress as large as 3.3 GPa at the surface. However, after yielding, the residual stress was released and a negligible amount of stress remained.
before yielding after yielding
0.3780 stress= -3.3GPa 0.3775
0.3770 stress= 0.1GPa
a (nm) 0.3765
0.3760
0.3755
0.00 0.05 0.10 0.15 0.20 0.25 sin2ψ
Figure A18: XRD measurement of surface residual stress magnitude
3. Hardness testing Vickers hardness indents were taken at randomly chosen positions in the gage area of two samples. One sample was in its as received condition and the other had been subjected to yielding. Table A7 lists the hardness tests results. The material’s hardness was calculated using the equation
…………………………………………………………..…..(7)
120
Table A7: Vickers hardness data obtained using 25 gram force As received sample hardness data After yielded sample hardness data (HV) (HV)
X1 X2 786 862 1066 929 1100 997 1010 833 783 940 739 944 1067 968 1054 801 1117 948
Using statistic methods, X1 = 969, S1=153; X 2 = 920 , S2=66. It is a small size sample, so student t test is proper to used to analyze these two groups of data. We can lookup in the table that for a confidence range of 95% and eight degrees of freedom, the value is 2.306. Therefore, calculated a 95% confidence interval for without yield is 844 854 Conclusion 1. Carburization process introduces a significantly high residual compressive stress on the alloy A286 material surface. 2. This surface stress was eventually completely released by yielding. 3. However, after eliminating compressive residual stress, there is no change in hardness that is statistically significant at the 95% confidence level. 121 Bibliography [1] Mechanical Behavior of Materials, N.E. Dowling, Prentice Hall, Upper Saddle River, NJ, 3nd Edition, 2007 [2] Fatigue of Materials, S. Suresh, 2nd edition, 1998 [3] F. Khoshnaw, A. Kheder, F. 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