effect of creep degradation on hardness changes of ni

ISIJ International, Vol. 59 (2019), No. 9

© 2019 ISIJ1695

ISIJ International, Vol. 59 (2019), No. 9, pp. 1695–1704

* Corresponding author: E-mail: [email protected]: https://doi.org/10.2355/isijinternational.ISIJINT-2018-878

1. Introduction

From the viewpoint of environmental protection, a policy of reducing CO2 emissions and fuel consumption based on the development of high efficiency thermal power plants was proposed by the Ministry of Economy, Trade and Industry of Japan.1) In the coal-fired power plants, the devel-opment of A-USC (Advanced Ultra-Supercritical) power technology, which raises the steam temperature up to 700°C from 600°C currently adopted, is considered as an important technology to increase the thermal efficiency. The basic design structure of the A-USC power boiler is as similar to the current USC boiler and thus it is possible to use most of their components to currently operated USC power boil-ers except the high temperature component which operates at 700°C. Therefore, the aged coal fired power plants can be easily up graded to A-USC power boiler and thus it is rapidly increasing as an optimal highly-efficient technology.

The establishment of life assessment method is essential to operate A-USC power plants safely for a long time.

Effect of Creep Degradation on Hardness Changes of Ni-based Alloys for A-USC Power Boiler

Nobuhiko SAITO,1)* Kazuki ISHIYAMA,2) Tomiko YAMAGUCHI3) and Fujimitsu MASUYAMA3)

1) Research & Innovation Center, Mitsubishi Heavy Industries, Ltd., 5-717-1 Fukahori-machi, Nagasaki, 851-0392 Japan2) Graduate Student at Department of Materials and Engineering, Kyushu Institute of Technology. Now at Nisshin Steel Co. Ltd, 1 Hama-machi, Hekinan, 447-0853 Japan.3) Department of Materials and Engineering, Kyushu Institute of Technology, 1-1 Sensui-cho, Tobata-ku, Kitakyushu, 804-8550 Japan.

(Received on December 28, 2018; accepted on March 11, 2019)

A project to develop Advanced Ultra-Supercritical (A-USC) power plant which operate at 700°C has been under way in Japan with the intension to reduce the CO2 emission and improve the power generation efficiency. Ni-based alloys, UNS N06674 (ASME Code Case 2826, 47Ni-23Cr-23Fe-7W, HR6W) and UNS N06617 (52Ni-22Cr-13Co-9Mo, Alloy 617) are candidate materials used in the high temperature regions of the A-USC power boilers. In this study, effect of creep degradation on the hardness change of both Ni-based alloys was investigated in order to evaluate the potential of the hardness method for creep life assessment. The investigation results showed that the scattering of hardness values in the latter stage of the creep life was increased as a result of hardening and softening of the material due to the increase of dislocation density with creep deformation and coarsening of precipitates followed by the formation of creep cavities, respectively. This tendency shows a good correlation with the creep life fraction. Therefore, it could be an useful tool for creep life evaluation. In addition, it was found that the hardness variation of the HR6W can be expressed as a function of temperature, stress and creep life fraction and thus leads to postulate a regression equation with hardness by a multivariate analysis. The creep life fraction can be calculated by substituting the values of service temperature and stress, and measured hardness into the multivariate regression equation. Those findings suggested that the hardness method was useful for the creep life assessment of Ni-based alloys.

KEY WORDS: A-USC power boiler; Ni-based alloy; creep life assessment; hardness; multivariate analysis.

Although it is possible to apply with existing life assessment methods, a method with better accuracy and simplicity is required. In the past years, the Ni-based alloy has not been used in the actual power plant boilers and thus its deteriora-tion and damage modes have not been understood clearly. Recently, there are several candidate methods proposed for Ni-based alloy. For example, focusing on the relationship between creep strain and GROD (Grain Reference Ori-entation Deviation), which is one of the EBSD (Electron BackScatter Diffraction pattern) parameters in the grain,2,3) the changes in the number density of creep cavities4) and the changes in hardness5,6) have been studied. Especially, the hardness method, which was originally developed for strain-softening steel such as creep strength enhanced fer-ric steel,7) is a simple and non-destructive method. If it is possible to clarify the relationship between creep life and hardness changes in Ni-based alloy, the method would be an extremely effective technique. Therefore, in this work, a fundamental study, regarding the relationship between hard-ness changes and creep damage progress of Ni-based alloys, UNS N06674 (ASME Code Case 2826, 52Ni-23Cr-23Fe-7W, HR6W, hereafter called HR6W) and UNS N06617 (52Ni-22Cr-13Co-9Mo, Alloy 617, hereafter called Alloy


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617) was conducted in order to evaluate the capability of the hardness method as a creep life assessment.

2. Materials and Experimental Procedure

Table 1 shows the chemical compositions of test materi-als of Ni-based alloys, HR6W and Alloy 617 which are candidate materials for thick-walled piping in A-USC power boilers. Mean grain size of HR6W and Alloy 617 were 220 μm and 250 μm respectively. The grain size was measured by optical microscope observation and manual mean linear intercept methods in accordance with ASTM E 112. Final heat treatment of both alloys was solution annealing. In order to evaluate the precipitation hardening due to thermal aging and strain hardening during creep, aged specimens and creep-interrupted specimens were prepared. The aging temperatures were 700, 750 and 800°C, and the aging time was ranging from 300 h to 10 000 h. The short-time aging test of 0.1, 0.3, 1, 3, 10, 30 and 100 h at 700°C was also conducted to investigate the changes in hardness at initial stage of aging. The creep-interruption was taken place at the life fraction of 20, 40, 60 and 80% before rupture, and creep test temperatures were 700 and 800°C. The stress condi-tions were 120 MPa to 200 MPa and 60 MPa to 100 MPa at 700°C and 800°C, respectively for HR6W, and 200 MPa to 350 MPa and 90 MPa to 140 MPa at 700°C and 800°C, respectively for Alloy 617.

The 20 points of Vickers hardness of aged specimens

were measured at the center of the specimens with the load of 1 kgf for 15 s. For the creep specimens, the Vickers hardness traverse from the head portion to the gauge portion (i.e. the center of the longitudinal direction or the near the fracture part) on the longitudinal section was measured at interval of 1 mm with the load of 1 kgf for 15 s. Measured points of each longitudinal location on the gauge portion, shoulder portion and head portion were 3 to 5, 5 to 6 and 7 respectively. In addition, the micro-hardness measurement was carried out on the creep specimens of HR6W tested at 700°C under 120 MPa and Alloy 617 tested at 700°C under 240 MPa with the load of 10 gf for 15 s to study the hardness distribution in the grains and its correlation to the crystal orientation which was analyzed by EBSD method after micro-hardness measurement. The 77 points of micro-hardness were measured at interval of 50 μm. The specimens used for EBSD analysis was polished using progressively finer grades of diamond paste, ranging from 3 μm to 1 μm in particle size, and then they were lightly etched using colloidal silica to finish it to a smooth and flat surface. The measurement by EBSD was carried out at an acceleration voltage of 20 kV, and the step size between unit pixels was 2 μm.

3. Results and Discussion

3.1. Effect of Thermal Aging on Hardness ChangesFigures 1 and 2 show the relationship between hardness

Table 1. Chemical compositions of test materials.

(mass%)

Alloy C Si Mn Ni Cr Mo W Co Ti Nb Al Fe

HR6W 0.08 0.4 1.2 43.0 23.0 – 6.0 – 0.08 0.18 – 26.1

Alloy617 0.07 0.48 0.52 52.1 22.1 9.0 – 12.7 0.32 – 1.19 1.52

Fig. 1. Relationship between hardness and aging time for HR6W. (Online version in color.)

Fig. 2. Relationship between hardness and aging time for Alloy 617. (Online version in color.)


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and aging time of HR6W and Alloy 617, respectively. In the case of HR6W, hardness after 300 h aging significantly increased to 200 HV from the initial hardness of 160 HV at every aging temperatures, and then gradually rose to a range from 220 to 240 HV up to 10 000 h. There was no significant difference in hardness values between the aging temperatures of 700, 750 and 800°C. On the other hand, in the case of Alloy 617, the hardness after 300 h aging increased to 240, 260 and 280 HV at 700, 750 and 800°C respectively from the initial hardness of 185 HV, and then the hardness remains constant at every aging temperatures. Saturated hardness at 700, 750 and 800°C are inferred to correlate with the equilibrium amount of γ ’ phase.8) The hardness of the materials aged for short time (0.1 to 100 h) at 700°C is shown in the graphs inserted in Figs. 1 and 2. The aging time on the horizontal axis is shown on a logarith-mic scale. The hardness of HR6W began to increase from 10 h, saturated once at 300 h, and then increased again after 3 000 h. As reported by Semba et al.,9) M23C6 precipitates in the short-time region and Laves phases continuously precipitates for a long time. Therefore, it is considered that the increase in hardness from 10 to 300 h and after 3 000 h aging corresponds to the precipitation of M23C6 and Laves phases respectively. The hardness of the Alloy 617 continu-ously increased even during the very early stage of aging due to the precipitation of γ ’ phase and reached saturation at about 300 h.

3.2. Effect of Creep on Hardness ChangesFigures 3 and 4 respectively show hardness distribution

from the head portion to the gauge portion in the longitudi-nal section of the creep specimens of HR6W and Alloy 617 interrupted at creep life fraction of 0.2, 0.4, 0.6 and 0.8. In the case of HR6W, hardness of the gauge portion was higher than that of the head portion due to the strain-hardening with the creep deformation. The hardness in the shoulder

portion rose linearly with increase in stress, and its slope became steeper with the increase of the creep life fraction. From these results, it was found that hardness change of tested Ni-based alloys clearly showed stress dependence. The hardness at head portion and gauge portion rose with increase in creep life fraction, and the hardness scatter in the gauge portion became broader with the creep life fraction. In the case of Alloy 617, hardness change with the increase in creep life fraction was small, and the change of hardness scatter was also small. It was reported that HR6W demon-strates superior rupture ductility, compared with Alloy 617

Fig. 3. Hardness distribution changes on the longitudinal section of creep specimens of HR6W. (Online version in color.)

Fig. 4. Hardness distribution on the longitudinal section of creep specimens of Alloy 617 (700°C, 240 MPa, tR=2 217.5 h). (Online version in color.)


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and other Ni-based alloys strengthened by γ ’ phase even in the long-term tests.9) Therefore, the strain hardening in HR6W became larger than Alloy 617 due to the difference of introduced creep strain.Figures 5 and 6 show the micro-hardness distribution

map in the grains near the creep cavities and IPF map super-imposed on IQ map obtained by means of EBSD analysis

to demonstrate the crystal orientation for creep rupture specimens of HR6W and Alloy 617. The hardness distribu-tion map is correlating with micro-hardness value in the IQ/IPF map. The numerical values in the hardness distribution map are micro-hardness values of each measurement point in the IQ/IPF map. And, the black lines and red lines show the boundaries of the grains and the cavities respectively.

Fig. 5. Micro hardness distribution in the grains and IQ/IPF map for HR6W crept at 700°C under 120 MPa.

Fig. 6. Micro hardness distribution in the grains and IQ/IPF map for Alloy 617 crept at 700°C under 240 MPa.


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The numerical values in the IQ/IPF map are the average micro-hardness values of each grain. The hardness in grain of both alloys shows large scatter, which seems to be caused by heterogeneous distribution of precipitates and pile-up of dislocations at the grain boundaries. The creep cavities were observed on the grain boundaries and the cavities preferentially generated at the boundaries between relatively hard grains and soft grains. In addition, the hardness of the crystal plane that near to (111) was also high.

3.3. Hardness Changes due to Creep DamageFigure 7 shows the hardness change at gauge portion

and head portion of creep-interrupted and ruptured speci-mens. Although the hardness at gauge portions and head portions of HR6W was increased with creep life fraction, the increasing tendency of hardness of HR6W at the gauge portions was greater than that of the head portions. On the other hand, the hardness at the gauge portions and the head portions of Alloy 617 was almost the same, except of

that under the test conditions of 700°C and 350 MPa. It is considered that hardness change of Alloy 617 with the low creep ductility strongly depend on the age hardening rather than creep strain hardening.Figure 8 shows the hardness change at gauge portion and

head portion of creep-interrupted and ruptured specimens of HR6W tested at 700°C under 120 MPa and Alloy 617 tested at 700°C under 240 MPa. The standard deviation of the hardness of all the specimens is shown as a band. It is found that the scatter band of the hardness of the HR6W with the large creep strain hardening became broad with increase in creep life fraction, which seems to be caused by heterogeneous distribution of local creep strain.Figure 9 shows the relationship between the hardness

increase due to creep (i.e. difference between the gauge-portion’s hardness and the head-portion’s hardness) and the hardness ratio (i.e. ratio of the gauge portion’s hardness to the head portion’s hardness), and the creep life fraction. With respect to the HR6W, both values rose with creep life

Fig. 7. Hardness changes at gauge portion and head portion of creep-interrupted and ruptured specimens. (Online ver-sion in color.)

Fig. 8. Hardness changes at gauge portion and head portion of the longest creep-interrupted and ruptured specimens. (Online version in color.)


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fraction except of the value at life fraction of 1.0, which had a large variation. As for Alloy 617, the changes of both values were small, but positive correlation was found in the test at 700°C under 350 MPa. Therefore, these parameters can be a candidate index to evaluate the creep damage.Figure 10 shows the relationship between the hardness-

increase (the data of ruptured specimen was excluded) and the creep strain of HR6W and Alloy 617. The relationship between hardness-increase and creep strain differed with each test condition, the change of hardness-increase in the both alloys at 700°C was larger than that of specimens ruptured at 800°C. If the creep strain can be estimated from the changes of hardness-increase, there is a possibility that the creep damage can be evaluated from the Ω method10,11) and Monkman-Grant’s law12) using the creep strain. Since the hardness-increase due to creep is corresponding to the dislocation density introduced by creep, it can be estimated by subtracting the hardness of the aged material (precipita-tion hardening) from the hardness measured in an actual component. In order to practically apply this method, it is necessary to establish a database of hardness change with respect to the precipitation hardening by conducting further hardness measurement test for actually serviced component

materials.Figures 11 and 12 show the frequency distribution of

hardness in terms of creep life fraction of HR6W tested at 700°C under 120 MPa and Alloy 617 tested at 700°C under 240 MPa. The number of measurements was 60 points or more on the gauge portion and 20 points or more on the head portion. In the case of HR6W, the difference between the hardness of the gauge portion and the hardness of the head portion spreads with increase in creep life fraction, except of the ruptured specimen. On the other hand, the hardness of Alloy 617 did not differ much between the gauge portion and the head portion at any life fraction. Fig-ure 13 shows the results from the analysis of the frequency distribution of hardness of HR6W and Alloy 617 shown in Figs. 11 and 12. It indicated that the distribution of hardness values in the gauge portion of HR6W spread with increase in creep life fraction.Figures 14 and 15 show the relationship between the

variance of hardness distribution and the creep life fraction in the gauge portion of each creep specimen tested at 700°C and 800°C. As it can be seen, at 700°C, the variance of Alloy 617 was found to be high at the early stage of life fraction, and then decreased once at t/tR of 0.7, and finally increased

Fig. 10. Relationship between hardness-increase due to creep and creep strain. (Online version in color.)

Fig. 9. Hardness-increase and hardness ratio at gauge portion and head portion of creep-interrupted and ruptured specimens. (Online version in color.)


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from the point exceeding t/tR of 0.7. On the other hand, the variance of HR6W increased with the creep life fraction. It is considered that the variance in the latter stage of life fraction is presumed to correspond to the hardness increase due to the increase of dislocation density accompanied with creep deformation and softening due to the coarsening of precipitates and the formation of creep cavities. Meanwhile, for 800°C, the variance of both alloys increased remarkably in the latter stage of life fraction. The reason for exceptional change of variance at 800°C is thought to be the accelera-tion of microstructural deterioration such as coarsening and

aggregation of precipitates. From the above results, there is a possibility that the creep life could be evaluated by means of hardness measurement of multiple points, calculating the variance of hardness data by statistical processing, and by associating the variance with the creep life fraction.Figure 16 shows the relationship between the hardness-

increase due to creep and Larson-Miller parameter. If the hardness-increase is a function of temperature and time, it should be uniquely expressed by single line. However, the relationship between the hardness-increase and Larson-Miller parameter showed a series of parallel line for each

Fig. 11. Frequency distribution of hardness in terms of creep life fraction of HR6W tested at 700°C under 120 MPa.

Fig. 12. Frequency distribution of hardness in terms of creep life fraction of Alloy 617 tested at 700°C under 240 MPa.


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test condition of stress. According to this result, it is clear that the hardness change of the Ni-based alloy cannot be represented uniquely by temperature and time only. Since the Monkman-Grant relationship and Norton’s law are expressed by Eqs. (1) and (2) respectively, the rupture time can be derived as Eq. (3). To introduce this equation into the term of time in Larson-Miller parameter, the Eq. (4) can be expressed by the function of temperature and stress as

Fig. 13. Changes in frequency distribution of hardness in terms of creep life fraction. (Online version in color.)

the Φ parameter.

�mtR MGK� ................................ (1)

� �m

n� A .................................. (2)

t nR MGK A� / � ............................. (3)

Fig. 14. Relationship between variance of hardness distribution and creep life fraction at 700°C.

Fig. 15. Relationship between variance of hardness distribution and creep life fraction at 800°C.


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Fig. 16. Relationship between hardness-increase due to creep and Larson-Miller parameter. (Online version in color.)

Fig. 17. Relationship between hardness-increase due to creep and φ parameter.

Fig. 18. Relationship between predicted hardness and measured hardness based on multivariate analysis using all data. (Online version in color.)

� � � � ��

��T C Log K AMG / � n ................. (4)

where εm is the minimum creep rate, tR is the time to rupture, KMG is the Monkman-Grant constant, n is the stress expo-nent, T is the absolute temperature and A and C are respec-tively the coefficient and Larson-Miller parameter constant.Figure 17 shows the relationship between the hardness-

increase with creep and Φ parameter of HR6W. In this sec-tion, only HR6W which could be calculated the Monkman-Grant relationship and Norton’s law based on sufficient creep data was evaluated. The Monkman-Grant constant (KMG) was estimated from the creep rupture test result of HR6W, and the material coefficient (A) and the stress expo-nent (n) of Norton’s law was obtained from the literature reported by Noguchi et al.13) Because the hardness-increase decreased linearly with the increase of Φ parameter, a nega-tive correlation was found between the hardness-increase and the Φ parameter. From the results, it is found that the reduction of the hardness-increase is occurred in the high temperature and the low stress conditions. In this way, by using the Φ parameter, the hardness-increase can be uniquely expressed by a single line as a function of tem-perature and stress. Therefore, the creep induced hardness-increase at rupture time can be estimated from the creep temperature and stress conditions.

Because the hardness changes of the HR6W have a cor-relation with temperature, stress and time (or life fraction), the hardness changes with creep should be expressed by the function of the Eq. (5).

H T� � �� f t or t tR, , /� ...................... (5)

where H is hardness, T is the temperature, σ is the stress, t is the time and t/tR is creep life fraction.

The multivariate analysis was carried out using all hard-ness data of gauge portion, head portion and shoulder por-tion in creep-interrupted and ruptured specimens of HR6W. The stress at the shoulder portion was calculated from the diameter at the hardness measurement position, and the stress at the head portion assumed zero. Figure 18 shows the relationship between measured hardness and predicted hardness based on multivariate analysis using all data. The scatter of the predicted hardness against the measured hard-

ness was extremely large, and good correlation cloud not be obtained between the both values. In the creep rupture specimens, because the hardness scatter was increased by


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the hardening due to the increase of dislocation density associated with creep deformation and the softening due to coarsening of precipitates and generation of creep cavi-ties, the variance of analysis results also is presumed to be became large. Therefore, re-multivariate analysis using all data excluding the data of the ruptured specimens was carried out. The calculation result is shown in Fig. 19. Although the scatter of the predicted hardness against the measured hardness was still large, a positive correlation was found between the predicted hardness and the measured hardness. The prediction accuracy of measured hardness at 99% upper and lower confidence limit was ±18.5% (±18.5 HV). Because the service temperature and stress are known in the actual component of thick-walled piping, the creep life fraction can be calculated by substituting the service temperature and stress, and measured hardness into the regression equation calculated by multivariate analysis. Therefore, it is considered that this method is useful for the creep life assessment. In the future study, it is necessary to collect a large number of data under different creep rupture test conditions and to improve prediction accuracy by opti-mization of multivariate regression equation.

4. Conclusion

The effect of creep degradation on hardness changes of Ni-based alloys, HR6W and Alloy 617 were investigated in order to study the capability of hardness method as a creep life assessment technique. As a result, the following conclu-sions were obtained.

(1) In the hardness changes due to thermal aging, the hardness of HR6W increased from 10 h to 300 h correspond-ing to the precipitation of M23C6, and then the hardness increase saturated once at 300 h, and increased again after 3 000 h up to 10 000 h corresponding to the precipitation of Laves phases. There was no difference in hardness between the aging temperatures of 700, 750 and 800°C. On the other hand, the hardness of Alloy 617 continuously increases from

the very early stage of aging due to the precipitation of γ ’phase and its hardening are saturated at about 300 h with different hardness levels at each aging temperature.

(2) The hardness distribution in the longitudinal section of the creep specimens showed clearly stress dependence, and the hardness scatter increases with the progress of creep damage. It was presumed that the large scatter in the latter stage of creep caused by the simultaneous occurrence of hardening due to the increase of dislocation density with creep deformation and softening due to the coarsening of precipitates and the formation of creep cavities. The creep cavities preferentially generated at the boundaries between relatively hard grains and soft grains.

(3) The hardness-increase due to creep and the hardness ratio of the load portion to the unload portion tends to rise with increase in creep life fraction. These tendencies were particularly large in HR6W. Therefore, these parameters can be a candidate index to evaluate the creep damage.

(4) The variance of hardness shows a good correlation with the creep life fraction. Therefore, there is a possibility that the creep life could be evaluated by means of hardness measurement of multiple points, calculating the variance of hardness by statistical processing, and by associating the variance with the creep life fraction.

(5) The regression equation for hardness of HR6W, which consists a function of temperature, stress and creep life fraction, can be obtained by the multivariate analysis. Because the service temperature and stress are known in the actual component of thick-walled piping, the creep life frac-tion can be calculated by substituting the service tempera-ture and stress, and measured hardness into the regression equation. Therefore, this method is useful for the creep life assessment. In the future study, it is necessary to collect a large number of data under different creep rupture test con-ditions and to improve prediction accuracy by optimization of multivariate regression equation.

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Fig. 19. Relationship between predicted hardness and measured hardness based on multivariate analysis using data excluding the data from ruptured specimens. (Online version in color.)

effect of creep degradation on hardness changes of ni

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