1

Oxidation Behaviors of an Alloy617 in Very High Temperature Air and

Helium Environments

Changheui Jang*, Daejin Lee, and Daejong Kim

Department of Nuclear and Quantum Engineering, Korea Advanced Institute of Science and Technology,

373-1 Guseong-dong Yuseong-gu, Daejeon, 305-701, Republic of Korea,

*Tel.: +82-42-869-3824, Fax: +82-42-869-3810, e-mail: chjang@kaist.ac.kr

Abstract

The oxidation characteristics of Alloy617, a candidate structural material for the key components in the

very high temperature gas-cooled reactor (VHTR) application, were investigated. High temperature oxidation

tests were conducted at 900oC and 1100oC in air and helium environments and the results were analyzed.

Alloy617 showed parabolic oxidation behaviors at 900oC, but unstable oxidation behaviors at 1100oC, even in a

low oxygen containing helium environment. The SEM micrographs observation also revealed that the surface

oxides became unstable and non-continuous as the temperature or the exposure time increased. According to the

elemental analysis, the Cr-rich oxides were formed on the surface and the Al-rich discrete internal oxides were

formed below the surface oxide layer. After 100 hrs in 1100oC air, the Cr-rich surface oxide became unstable and

non-continuous, and the matrix elements like Ni and Co were exposed and oxidized. Depletion of grain

boundary carbides as well as matrix carbides was observed during the oxidation in both environments. When

tensile loading was applied during the high temperature oxidation, the thickness of the surface oxide layer, the

internal oxidation, and decarburization were enhanced because of the increase in diffusion of oxidizing agent and

gaseous reaction products. Such enhancement would have detrimental effects on the high temperature

mechanical properties, especially the creep resistance of Alloy617 for the VHTR application.

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1. INTRODUCTION

The very high temperature gas-cooled reactor (VHTR) is currently the most promising reactor type among

the generation-IV (Gen-IV) reactors for producing electricity and hydrogen economically. It is a graphite-

moderated reactor which uses helium gas to remove the heat generated in reactor core to the power conversion

unit or the hydrogen production plant. To achieve the high thermal efficiency and the hydrogen production

capability, the helium coolant will be operated at the temperature of about 850oC or higher [1]. Also, the coolant

gas pressure of up to 8 MPa is being considered. Therefore, the structural materials of the VHTR have to survive

such a harsh condition for the lifetime of the reactor, or up to 60 years. In the VHTR, the high temperature

helium coolant from the reactor core will go through the hot-gas duct (HGD) and the intermediate heat

exchanger (IHX) where the heat is transferred to the power conversion system and the hydrogen production

facility. Among the variety of materials reviewed, several nickel-base superalloys are seriously considered as the

materials for these high temperature components [2].

Nickel-base superalloys, especially the precipitation hardening superalloys have been widely used in many

high temperature applications because of high temperature strength, oxidation resistance, creep resistance, and

phase stability [2]. However, the proposed operating temperature of the VHTR is so high that the precipitates

responsible for the high temperature strength of those alloys are no longer stable. Because of the dissolution of

the precipitates, the precipitation hardening superalloys are not suitable for the VHTR application. On the other

hand, the solid solution hardening superalloys which depend on the lattice distortion caused by alloying elements

for their strength, can maintain considerable strength at higher temperature [2-4], and considered suitable for the

VHTR application. Among these superalloys are Alloy617, Haynes230, and Hastelloy-XR.

When superalloys are exposed to the VHTR coolant for prolonged period, they will be subjected to the

degradation due to creep [2,5]. Therefore, the creep properties of the candidate superalloys should be well

understood and characterized to be used in design and analysis. In addition, oxidation or other high temperature

corrosion resistances of these superalloys are of concern because they would have effects on the mechanical

property degradation of superalloys by metal loss, decarburization, internal oxidation, etc. [5-7]. Therefore, to

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understand creep resistance of the candidate superalloys, the oxidation behavior and the stability of oxides

should be properly understood. Accordingly, there have been several studies on the high temperature oxidation of

superalloys in various helium environments [7-16]. However, the evolution of oxides and microstructure changes

of Alloy617 during the high temperature oxidation have not been clearly identified. Furthermore, the effects of

tensile loading on the oxidation behavior of Alloy617 have not been properly investigated.

Therefore, in the present work, the oxidation behaviors of Alloy617, a one of the candidate alloys, were

investigated by oxidation tests at 900oC and 1100oC in air and impure helium environments for up to 200 hrs.

The evolution of oxides and microstructure changes during the high-temperature oxidation were also analyzed

through the scanning electron microscope observation and the elemental analysis of the oxidation tested

specimens. Finally, the effects of tensile loading on the oxidation behaviors and microstructure changes were

investigated by analyzing the creep specimens tested in high temperature air.

2. EXPERIMENTAL PROCEDURES

2.1. Test material and specimens

The material used in this study is an 18 mm-thick plate of Inconel Alloy617 produced by Special Metals

Corporation. Chemical composition of the material is shown in Table 1. The material was solution annealed at

1175oC for an hour and water quenched. The typical microstructure is shown in Fig. 1. As shown in the figure,

the as-received microstructure contains twins, well-distributed primary carbides and grain boundary carbides.

These carbides were known to be primarily M23C6 type carbides precipitated during solution anneal heat

treatment in Alloy617 [8]. According to the supplier’s information [12], the average ASTM grain size is 4.6. The

mechanical properties at room temperature are summarized in Table 2. As shown in Fig. 2, the oxidation test

specimen used in this study is a 1 mm-thick coupon type specimen. The round bar type specimen shown in the

figure was used in creep tests and analyzed to investigate the effects of tensile loading on the oxidation behaviors.

<Table 1>

4

<Table 2>

<Fig. 1>

<Fig. 2>

2.2. Test setup and conditions

The oxidation tests for coupon type specimens were conducted in air and helium gas environment at 900oC

and 1100oC. A box furnace was used for the oxidation tests in air. The specimens were exposed at the test

conditions for up to 100 hrs. In helium test, vacuum furnace shown in Fig. 3 was used. The helium gas used in

this study is of 99.999% high purity containing small amount of impurities, such as 1.8 ppm H2O, 1.4 ppm O2,

3.1 ppm N2. Before the tests, the vacuum furnace was purged with helium after achieving 10-3 torr vacuum. This

process was repeated 3 times. Then, a constant helium flow rate at 60 cc/min was maintained for 5 hours at room

temperature. Finally, the furnace was heated up to the test temperature in about 2 hrs. The helium test was

conducted up to 200 hrs with the helium flow rate at 60 cc/min. The creep tests using the round bar type

specimens were conducted in air at 900oC using a servo-electric loading machine equipped with an induction

heater.

<Fig. 3>

2.3. Measurement of weight change and microstructure analysis

To measure the weight change due to oxidation, a specimen was removed from the furnace at each pre-

determined hour. To measure the weight, a microbalance with accuracy of ±0.00005g was used. Then, the

oxidation rate was calculated as the weight gain per unit surface area, ∆W/S (g/m2), from the measured weight

change using following equation:

oot SWWSW /)(/ −=∆ (1)

where, Wo is the weight before oxidation, Wt is the weight after oxidation, and So is the surface area before

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oxidation. After the measurement of the weight gain, specimens were prepared for the microstructure

observation. The specimens were cut normal to the surface and polished, etched at room temperature in 100 ml

HCl/0.5g-CrO3 solution for 60 to 80 sec, and gold-coated. Then the surface and sub-surface oxides and

microstructure were observed under the scanning electron microscope (SEM). For composition analysis of the

elemental mapping, an energy dispersive X-ray analysis (EDX) was used.

3. RESULTS AND DISCUSSION

3. 1. Weight gain in high temperature air and helium environments

The weight gain test results are summarized in Fig. 4. From the figure, a similar weight gain behavior was

observed at 900oC in both air and helium environments despite of the huge difference in oxygen content between

air (200000ppm) and helium (1.4ppm). Also, the weight gain followed a parabolic curve at 900oC. However, the

weight gain behaviors at 1100oC were quite different from those at 900oC. At 1100oC, after initial weight gain

behaviors following a parabolic curve, decrease in weight gain, or the turnaround, was observed, which resulted

in weight loss at longer oxidation time. Also, the turn-around of the weight gain happened later and the

maximum weight gain at the turnaround was greater in helium environment than in air condition.

<Fig. 4>

The potential sources of weight change during high temperature corrosion tests are oxidation, spallation,

carburization, decarburization, and etc. [13]. However, in our tests, carburization was not expected to happen

because there is no carbon containing impurities in helium gas. In case of Alloy617, oxidation of alloying

elements like Cr, Ni, Al, and Ti at the metal surface and inner matrix could be the reason of weight gain. On the

other hand, the sources of weight loss are spallation of oxides, evaporation of oxides, decarburization, and

thermal etching [5-7,14]. The proofs of the spallation, or tiny black particles were observed on the specimen

holders around the vertically placed specimens in 1100oC tests, but not in 900oC tests. In addition to the

spallation, another major source of weight loss is the evaporation of Cr-rich oxide layers which become

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thermodynamically unstable and begin to be evaporated at 1100oC [17]. Other sources of weight loss, like

decarburization and thermal etching were also observed, but their contribution to the weight loss was considered

not so significant. Therefore, the sharp drop of weight gain in 1100oC was thought to be primarily caused by the

spallation and vaporization of Cr-oxides during the extended exposure. In high temperature helium containing

variety of impurities, Christ et al. [16] reported a stable weight gain behavior at 900oC for up to 150 hrs.

However, at 1000oC, they observed initial weight loss caused by decarburization prior to stable weight gain

afterward. On the other hand, Ganesan et al. [9] reported a similar behavior as observed in our study, though the

turnaround time was about 3 times larger in air + 10% water vapor at 1100oC. Therefore, it could be said the Cr-

rich surface oxides would be stable in oxidizing environments at up to 1000oC, but became unstable at 1100oC.

Based on the weight gain test results, it was found that Alloy617 did not have the proper oxidation resistance at

1100oC even in a very low oxygen containing environment.

3.2. The morphology and stability of the oxide layer

The SEM micrographs of specimens oxidized for 100 hrs are shown in Fig. 5 and 6. In general, similar

morphologies were observed in air and helium environments. Thin outer oxide layer and inwardly protruded

oxides, or the internal oxidation were present. Also present was the area free of grain boundary carbides, or the

decarburized zone, just below the surface oxide layer. However, the thickness of oxides, the penetration depth of

internal oxidation, and decarburized zone were smaller in helium environments than in air at both test

temperatures as summarized in Table 3. At 900oC, the depth of the decarburized zone was similar in air and

helium environments. However, the surface oxide in air showed a relatively rough surface while that in helium

looked smooth and continuous as shown in Fig. 5. Thus, the larger penetration depth of the internal oxidation in

air could have been caused by the enhanced diffusion of oxygen through the less dense surface oxides in air. The

gaseous product of carbide decomposition reactions [13-16] could easily escape through the oxides to the

environment, too. Therefore, at 900oC, it could be said that the surface oxide was less dense and the oxidation

resistance was less in air than in helium environment despite of the similar weight gain behaviors shown in Fig.

7

4.

<Fig. 5>

<Fig. 6>

<Table 3>

As shown in Fig. 6 and Table 3, oxidation exposure for 100 hrs at 1100oC resulted in further growth in oxide

layers, internal oxidation and decarburized zone in both air and helium environments. On the other hand, a rather

thin and continuous surface oxide layers were observed in specimens exposed for 5 hrs in 1100oC air and helium

environments, which was in accordance with the observed weight gain shown in Fig. 4. However, at later hours,

the observed microstructure alone could not explain the weight gain behavior because the micrographs shown in

Fig. 6 were the combined results of several competing phenomena happened simultaneously. The thicker oxide

layer and deeper penetration of internal oxidation shown in the figures would have contributed to the weight gain

because the effect of decarburization on weight loss was not so significant. However, as shown in Fig. 4, the

weight gain was decreasing at 100 hrs in both 1100oC air and helium environments, implying that other

phenomena responsible for the weight loss were active. The other phenomena would be the spallation or the

evaporation of the surface oxides. As previously mentioned, the proof of spallation was observed on the

specimen holders. Even though it was not easy to observe the proof of oxide evaporation in this study, previous

study [18] and the rather discontinuous and scattered oxide layer shown in Fig. 6 could be the results of such

phenomena.

The SEM micrographs observation revealed that a continuous outer oxide layer and relatively less internal

oxidation were formed at 900oC tests than at 1100oC tests. It was considered that the continuous outer oxide

layer acted as barriers to the oxygen diffusion to the metal/oxide interface and slowed down the further oxidation

at 900oC. However, the micrographs of specimens tested at 1100oC showed discontinuous and scattered surface

oxide layers and exposed matrix metal at certain locations. These discontinuous oxide layers were not protective

enough to slow down the diffusion and migration of oxygen [16,18]. Therefore, a significant internal oxidation

along with a significant metal loss by oxidation and spallation occurred at 1100oC.

The detailed SEM micrographs of oxidized specimens tested in air for 100 hrs are shown in Fig. 7. The

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formation of voids within the oxides and breakdown of the surface oxide layers are evident in the figure. It is

clear in Fig. 7 (b) that after 100 hrs at 1100oC, the surface oxides were so unstable that some of them were about

to be detached from the surface revealing the matrix below. The detachment of the surface oxide, or the

spallation would have contributed the decrease in weight gain as shown in Fig. 4. Also shown in Fig. 7 is the

extensive formation of internal oxide below the unstable surface oxide layer. The composition and formation

mechanism of the internal oxides will be discussed later.

3. 3. Evolution of oxide layers and internal oxidation

The EDX analysis results of the surface oxide and the microstructure below are shown in Fig. 8. The

elemental mapping was done for the specimens tested for 5 hrs and 100 hrs at 900oC and 1100oC in air

environment. At 900oC, the surface oxide was mostly Cr-rich oxide, or Cr2O3 when the exposure time is 5 hrs.

Furthermore, Al was slightly enriched just below the surface oxide layer, forming semi-continuous Al-oxides. No

internal oxidation was evident at this early stage of oxidation. At much later, or 100 hrs of exposure, the surface

oxide was still predominantly Cr2O3, but a little more partitioning of Ti and Mn was observed. The most

significant difference between Fig. 8 (a) and (b) is the existence of the discrete Al-rich areas below the surface

oxide after 100 hrs of oxidation test. The Al-rich particles were thought to be Al-oxides, or Al2O3.

The evolution of the surface oxides and internal oxides were similar at 1100oC. Compared to the 900oC

results, a slight partitioning of Ti was observed. However, the amount of Ti partitioning in Cr-oxide was far less

compared to the results exposed for 1000 hrs in HTGR helium condition [15] despite of the extensive Al-rich

oxide formation below the surface oxides. However, the Cr content in the surface oxide became non-uniform

after 100 hrs at 1100oC, as shown in Fig. 8 (d). Instead, Ni and Co were enriched in areas of lower Cr-content on

the surface. This means that the Cr-rich oxide layer was no longer stable or continuous, and the matrix elements

like Ni and Co were exposed and oxidized where the Cr-oxides were damaged. The formation of such oxide-free

surface was also reported in the highly decarburized specimens [16]

Al was added to nickel-base superalloys primarily improve the high temperature oxidation resistance by

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partitioning into and stabilizing the surface oxides like Cr2O3 [18]. However, no evidence of Al partitioning into

the surface oxide was observed in this study. Instead, Al formed discrete internal oxide in the matrix below the

surface oxides. The internal oxidation would have been formed by the diffusion of oxygen through the surface

oxide and matrix. Therefore, the internal oxidation was promoted and formed at much below the surface as the

temperature increased and the surface oxide became unstable.

<Fig. 8>

3.4. Decarburization

As shown previously, the grain boundary near the surface oxides became gradually depleted of the carbides

as the oxidation progressed. The depletion of grain boundary carbides would have been caused by the reaction

between grain boundary carbide and oxygen provided by diffusion from the environment or oxidation/reduction

reaction of surface oxide layers [13,14,16]. It is well known that the grain boundary carbide prevents grain

boundary migration and sliding, therefore optimal distribution of grain boundary carbide would improve the

creep resistance of high temperature alloys [5,14]. Because our test environments were oxidizing condition

without containing reducing agent like CH4 or CO, significant decarburization was expected and the results

shown in Fig. 5 and 6 confirmed our assumption. As shown in the figures and Table 3, the decarburized zone was

greater as the temperature and exposure time increased. However, the effect of the test environment on the

decarburized depth was greater at 1100oC than at 900oC. This could be attributed to the relative stability of the

surface oxide layer in helium, as shown in Fig. 4.

When the superalloys were exposed at high temperature environment containing oxidizing agent like O2

and H2O, the carbide exposed on the surface and along the grain boundary would react with oxidizing agents

[13,16] and gradually disappear as shown in Fig. 9. It is clear in the figure that the dispersed carbides within the

grains were mostly disappeared after oxidation tests in air and helium environments. Decarburization also occurs

far below the surface as shown in Fig. 5 and 6. However the degree of decarburization observed in this study is

far less than the previously reported results [16]. In that result, the 2 mm thick coupon was completed depleted

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with grain boundary carbides when exposed for 150 hrs in 1000oC in helium with H2-H2O impurity, while the

depth of the decarburized zone was about 230µm at 950oC. They also suggested that the increase in oxygen

content in the gas would lower the extent of the decarburization, because of the stabilization effect of the surface

oxides at higher oxygen content. However, in out tests, the surface oxides were not stable in a highly oxidizing

condition, and the decarburized depth was greater. Also, despite of the increased stability of the surface oxides in

helium, considerable decarburization was observed, especially at 1100oC. Further investigation is needed to

clarify the effects of the oxidizing agents on the degree of the decarburization of Alloy617 at the VHTR

temperatures.

<Fig. 9>

3. 5. Effects of tensile loading on oxidation behaviors

The microstructures of the creep tested specimen at 48 MPa in 900oC air are shown in Fig. 10. The

specimen was sliced along the loading direction to examine the effects of tensile loading on the microstructure of

Alloy617. As shown in Fig. 10, many crack-like features and voids were observed near the surface of the highly

strain region (Fig. 10 (a)). At the region further away from the surface, very fine and serrated grains along with

the voids were observed despite of the very high strain under the tensile loading condition (Fig. 10 (b)). The fine

and serrated grain was considered as the evidence of the dynamic recrystallization (DRX). DRX is the

strengthening mechanism during hot deformation caused by the generation and migration of high angle

boundaries inside the sub-grains [19,20]. At areas farther away from the rupture region, no DRX was observed

because the strain was not high enough.

<Fig. 10>

Fig. 11 shows the effect of tensile loading on the oxidation and decarburization of Alloy617. On the figure,

it is clear that the thickness of the surface oxide layer was thicker and the internal oxidation was deeper in the

loaded specimen than in the unloaded oxidation specimen. Similar behavior was also observed by Mino et al.

[10]. They suggested the enhanced intergranular oxidation was attributed to the localized fracture of external

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oxide layers and internal Al-oxides due to grain boundary sliding. Enhanced oxidation and decarburization were

clear in our tests as shown in Fig. 11 and 12 and summarized in Table 4. It is thought that these differences

would have been caused by the enhanced diffusion of oxidizing agent and gaseous reaction products formed

along the grain boundary in the tensile loaded specimen. As the creep resistance is affected by the grain

boundary carbides, the enhanced carbide depleted zone under the tensile loading would have detrimental effects

on the creep properties of the materials.

The various microstructure features of the tensile loaded specimen are shown in Fig. 12. It is clear from the

figures that the tensile loading caused surface cracking along the grain boundary, voids below the surface, and

voids within the matrix during the creep tests. It could be said that these features are closely related to the creep

resistance of the materials by providing the crack initiation and void nucleation sites during creep deformation.

Similar results on the contribution of the unstable surface oxides on the formation and growth of cracks during

creep tests in high temperature air were reported [14]. However, in out tests, the triple point carbides, or the

carbide particle gathered at triple point to release accumulated stress [14] was not observed because of the

relatively high stress level.

<Fig. 11>

<Fig. 12>

<Table 4>

4. CONCLUSIONS

High temperature oxidation tests were performed at 900oC and 1100oC in air and helium environments to

investigate the stability of the oxide and the change of the microstructure of Alloy617, a candidate superalloys

for the VHTR application. Additionally, the effects of tensile loading on oxidation behaviors and microstructure

were also evaluated analyzing the creep specimens tested in air. Based on the oxidation tests and subsequent

microstructure analyses, the following conclusions were drawn.

1. At 900oC, Alloy617 showed a similar weight gain behavior in both air and helium environments despite of

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the huge difference in oxygen content. However, at 1100oC, after initial weight gain, the surface oxides

became unstable and the weight loss was observed due to the spallation and evaporation of the surface

oxides, even in low oxygen containing helium environment.

2. The SEM micrographs observation revealed the presence of the thin outer oxide layer and the internal

oxidation, and the decarburized zone in oxidation tested specimens. Such features increased as the

temperature or the exposure time increased. In high temperature air, the formation of voids within the oxides

and breakdown of the surface oxide layers were so extensive that they could not provide proper oxidation

protection.

3. During the oxidation, the Cr-rich oxides were formed on the surface with a slight partitioning of Ti and Mn.

However, Al formed the discrete internal oxides below the Cr-rich surface oxides by the diffusion of oxygen

through the surface oxide and matrix. After 100 hrs in 1100oC air, the Cr-rich surface oxide was no longer

stable or continuous and the matrix elements like Ni and Co were exposed and oxidized. Also, the depletion

of grain boundary carbides was observed in air and helium environments, and the effect of the environment

was greater at 1100oC. The decarburized zone became greater as the temperature and exposure time

increased.

4. When the tensile loading was applied, the thickness of the surface oxide layer, the internal oxidation, and

decarburization were enhanced by the enhanced diffusion of oxidizing agent and gaseous reaction products

along the grain boundary. Along with the enhanced carbide depleted zone, various microstructure features

like surface cracking along the grain boundary, voids below the surface, and voids within the matrix would

have detrimental effects on the creep properties of the materials.

ACKNOWLEDGEMENTS

This study was supported by the Korean Ministry of Science and Technology (MOST) under the Basic

Atomic Energy Research Institute (BAERI) Program and the US-Korea i-NERI Program. Part of the funding

was provided by the Second Phase BK21 Program of the Ministry of Education and Human Resource

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Development of Korea.

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8. Markins W., Hosier J., Bassford T. Microstructure and phase stability of Inconel alloy 617. Metall Trans

1974;5:2579-2590.

9. Ganesan P., Smith G., Yates D. Performance of Inconel alloy 617 in actual and simulated gas turbine

environments. Materials and Manufacturing Processes 1995;10(5):925-938.

10. Mino K., Ohtomo A., Saiga Y. The effect of stressing on the intergranular oxidation of Inconel 617. J Japan

Inst Metals 1980;44(12):1397-1403.

11. Chin J., Johnson W., Chen K. Compatibility of aluminide-coated Hastelloy X and Inconel 617 in a simulated

gas-cooled reactor environment. Thin Solid Films 1982;95:85-97.

12. Special Metals Corporation. Inconel alloy617. Technical Bulletin, 2005.

13. Kaczorowski D., Chapovaloff J. Corrosion behaviour of Inconel 617 in VHTR environment. Proceedings of

ICONE14 2006; ICONE14-89014.

14. Cook R. Creep properties of Inconel-617 in air and helium at 800 to 1000 oC. Nuclear Technology

1984;66:283-288.

15. Hirano T., Okada M., Araki H., et al. Corrosion of Inconel 617 in HTGR grade helium. Metall Trans A

1981;12A:451-457.

16. Christ H., Künecke U., Meyer K., et al. High temperature corrosion of the nickel-based alloy Inconel 617 in

helium containing small amounts of impurities. Mat Sci and Eng 1987;87:161-168.

17. Seal S., Kuiry S., Bracho L. Studies on the surface chemistry of oxide films formed on IN-738LC superalloy

at elevated temperatures in dry air. Oxidation of metals 2001;56(5/6):583-603.

18. Huang J. High-temperature oxidation behavior and mechanism of a new type of wrought Ni-Fe-Cr-Al

superalloy up to 1300 oC. Oxidation of Metals 2000;53(3/4):273-287.

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19. Brünger E., Wang X., Gottstein G. Nucleation mechanisms of dynamic recrystallization in austenitic steel

alloy 800H. Scripta Materialia 1998;38(12):1843-1849.

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Table captions

Table 1. Chemical composition of Alloy617 used in the study

Table 2. Mechanical properties of Alloy617 at room temperature

Table 3. Comparison of oxide layer thickness, internal oxidation depth, and decarburization zone depth

depending on the oxidation test conditions

Table 4. Effects of the tensile loading on the oxide layer thickness, internal oxidation depth, and decarburization

zone depth in 900oC air environment

16

Figure captions

Fig. 1. Microstructure of Alloy617, as-received condition

Fig. 2. Oxidation test specimens, (a) Coupon type and (b) Round bar type

Fig. 3. Schematic diagram of oxidation test setup in helium environment

Fig. 4. Weight gain versus oxidation time in various test conditions

Fig. 5. SEM micrographs of oxidized specimens tested in air and helium for 100hrs at 900oC. (Bottom photos are

the enlarged view of the area in the square on the top photos)

Fig. 6. SEM micrographs of oxidized specimens tested in air and helium for 100hrs at 1100oC. (Bottom photos

are the enlarged view of the area in the square on the top photos)

Fig. 7. SEM micrographs of oxidized specimens tested in air for 100hrs at 900 and 1100oC

Fig. 8. EDX analysis result of oxidized specimens tested in air, (a) for 5hrs at 900oC, (b) for 100hrs at 900oC, (c)

for 5hrs at 1100oC, (d) for 100hrs at 1100oC

Fig. 9. SEM micrographs of the oxidized specimen tested in air and He, taken at the center of the specimens

Fig. 10. Optical microstructure near the rupture region of the creep tested specimen, under 48MP of tensile creep

loading at 900oC

Fig. 11. Comparison of the microstructures of loaded and unloaded specimens; (a) at uniform elongation area of

the creep tested specimen, 82hrs in air at 900oC and (b) oxidation tested specimen for 100hrs in air at

900oC.

Fig. 12. Various features observed in the creep tested specimen (stress = 60MPa, for 5hrs in air at 900oC,

diameter reduction = 0.09); (a) Near surface optical microstructure, (b) Inner matrix SEM microstructure,

and (c) Near surface SEM microstructure, compared with (d) oxidation tested specimen for 5hrs in 900oC

air.

17

Table1. Chemical composition of Alloy617 used in the study

element C Mn Fe S Si Cu Ni Cr Al Ti Co Mo P B

w/o .09 .06 1.16 .001 .08 .05 53.94 21.68 1.14 .52 11.53 9.74 .006 .002

Table 2. Mechanical properties of Alloy617 at room temperature

Yield strength Ultimate tensile

strength % elongation

Reduction of

area

Average grain

diameter

364 MPa 823 MPa 53.3 % 55.9 % 71.9 mm

Table 3. Comparison of oxide layer thickness, internal oxidation depth, and decarburization zone depth

depending on the oxidation test conditions

Temperature Environment Outer oxide layer

thickness, [µm]

Internal oxidation

depth, [µm]

Decarburization zone

depth, [µm]

900oC Air 4 10 25

Helium 3 7 20

1100oC Air 10 42 200

Helium 8 40 80

Table 4. Effects of the tensile loading on the oxide layer thickness, internal oxidation depth, and decarburization

zone depth in 900oC air environment

Loading condition Duration Outer oxide layer

thickness, [µm]

Internal oxidation

depth, [µm]

Decarburization

zone depth, [µm]

No 100hrs 4 10 25

48MPa creep load 82hrs 8 30 80

No 5hrs 2 2 10

60Mpa creep load 5hrs 6 25 50

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Fig. 1. Microstructure of Alloy617, as-received condition

(a) (b)

Fig. 2. Oxidation test specimens, (a) Coupon type and (b) Round bar type

Fig. 3. Schematic diagram of oxidation test setup in helium environment

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Fig. 4. Weight gain versus oxidation time in various test conditions

(a) Air at 900oC (b) He at 900oC

Fig. 5. SEM micrographs of oxidized specimens tested in air and helium for 100hrs at 900oC. (Bottom photos are

the enlarged view of the area in the square on the top photos)

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(a) Air at 1100oC (b) He at 1100oC

Fig. 6. SEM micrographs of oxidized specimens tested in air and helium for 100hrs at 1100oC. (Bottom photos

are the enlarged view of the area in the square on the top photos)

(a) for 100hrs at 900oC (b) for 100hrs at 1100oC

Fig. 7. SEM micrographs of oxidized specimens tested in air for 100hrs at 900 and 1100oC

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(a) for 5hrs at 900oC

(b) for 100hrs at 900oC

(c) for 5hrs at 1100oC

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(d) for 100hrs at 1100oC

Fig. 8. EDX analysis result of oxidized specimens tested in air, (a) for 5hrs at 900oC, (b) for 100hrs at 900oC, (c)

for 5hrs at 1100oC, (d) for 100hrs at 1100oC

Fig. 9. SEM micrographs of the oxidized specimen tested in air and He, taken at the center of the specimens

900oC, air, 5h 1100oC, air, 5h 900oC, He, 5h

1100oC, He, 100h 100h 100h 100h

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Rupture, 82hrs, 48MPa (a) Surface (b) Inside

Fig. 10. Optical microstructure near the rupture region of the creep tested specimen, under 48MP of tensile creep

loading at 900oC

Rupture, 82hrs , 48MPa (a) tensile loaded specimen (b) oxidation specimen

Fig. 11. Comparison of the microstructures of loaded and unloaded specimens; (a) at uniform elongation area of

the creep tested specimen, 82hrs in air at 900oC and (b) oxidation tested specimen for 100hrs in air at 900oC.

0

0

DDD −=0.42

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0

0

DDD −=0.06

Loading direction

8µm 30µm

80µm

25µm

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(a) near surface (b) inner matrix

(c) near surface (d) Oxidation tested specimen

Fig. 12. Various features observed in the creep tested specimen (stress = 60MPa, for 5hrs in air at 900oC,

diameter reduction = 0.09); (a) Near surface optical microstructure, (b) Inner matrix SEM microstructure, and (c)

Near surface SEM microstructure, compared with (d) oxidation tested specimen for 5hrs in 900oC air.

Loading direction

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20µm