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Development of high temperature structural materials for the HTGR
Nuclear Engineering and Design 144 (1993) 305-315 North-Holland
305
Development of high temperature structural materials for the HTGR Y. Muto, H. N a k a j i m a a n d M. Eto Tokai Research Establishment, Japan Atomic Energy Research Institute, 2-4 Shirane, Shirakata, Tokai-mura, lbaraki-ken, 319-11, Japan Received 24 May 1993
In accordance with the HTGR program in Japan, a series of R&D for high temperature structural materials in particular with respect to the HTTR design code has been performed in JAERI for more than 20 years. This paper introduces R&D results of the pressure retaining low alloy steel 2 1/4Cr-lMo and the high temperature structural alloys Hastelloy XR and Ni-Cr-W superalloy for the design code together with some fruits of recent studies.
1. Introduction Studies on metallic materials for high temperature gas-cooled reactor (HTGR) applications have been made at the Japan Atomic Energy Research Institute(JAERI) for more than 20 years. A high temperature design code for the High Temperature Engineering Test Reaetor(HTTR)[1] has been already organized based on the data obtained by a series of material testings and structural tests. The items of research activities carried out are classified into three categories, i.e., those for the pressure retaining low alloy steels, for the characterization of the high-temperature alloys and for their structural application at very high temperature. This paper presents an overview of the JAERI's program during the last decade and emphasizes recent substantial advancement in each category.
2. Development of alloys for the pressure boundary structural components for the HTrR Among various components in the nuclear reactor, the pressure vessel is the one which requires a very high reliability. During normal operation of the HTTR the temperature of the vessel is supposed to be 400°C. However, there would be a possibility that the temperature would rise to 440°C at an abnormal transient state so that the design temperature was set at this temperature. It was decided that 2 1 / 4 C r - l M o steel
would be employed for the pressure vessel of the HTTR in consideration of the higher service temperature comparing with the light water reactor [2]. R & D of the steel have been carried out focussing on the aging behavior in relation to the chemical composition, rolling and forging process, welding method, etc.. To characterize the material, tensile and fatigue strengths, fracture toughness, Charpy impact energy etc. have been measured for the irradiated and thermally aged specimens. Here, the results of characterization tests are summarized together with some fruits of recent studies. 2.1. Application of 2 1 / 4 Cr-lMo steel to the HTTR Comparison between the normalized and tempered (NT) material and the annealed (AN) one was made from various aspects. As a result, NT was selected on the basis of the following facts. (1) Both strength and toughness increase due to the NT treatment which also rises the design stress up to around 450°C. (2) Higher toughness is maintained at the lowest service temperature. (3) NT is more resistant to the radiation and thermal aging embrittlements. To specify the NT material for the pressure vessel the following procedures were taken on the basis of the data on strength, ductility and toughness of the thick plate [3]. (1) With regard to the thickness ( ~ 160 mm) of the plate, a quenched and tempered (QT) treatment equiv-
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Y Muto et al. / High temperature structural materials for HTGR
306
alent to NT was employed. The optimum post weld heat treatment (PWHT) parameter P was determined as
P = T(log t + 20) × 10 -3 = 2 0 . 5 -
20- 75,
(1)
where T and t are temperature and time in °C and hr, respectively. The effect of the normalizing treatment can be disregarded here because the NT temperature is set a little bit lower than usual so that the virtually PWHT is only an influential treatment. (2) For the reduction of thermal aging and radiation embrittlements, the requirements on the chemical composition were derived empirically, i.e., J = (Si + M n )( P + Sn) × 104 < 100,
(2)
10P + 5Sb + 4Sn + As =
100
< 10.
(3)
Here, Si, Mn, P, etc. are expressed in mass%. These requirements indicate that not only impurity elements such as Si and P but also alloying elements such as Mn and Sn are important in the determination of the embrittlement parameter.
2.2. Effect o f neutron irradiation embrittlement on the thermal aging embrittlement The neutron irradiation embrittlement of the NT material in the upper shelf region as well as the effect of neutron irradiation on thermal aging embrittlement was investigated in terms of elastic-plastic fracture toughness. The material used is shown in Table 1 [4]. Irradiated specimens were 4 mm diameter by 26 mm gauge length tensile specimens, standard size Charpy V-notch (CVN) specimens, and disk-shaped compact tension (DCT) specimens with three different thickness values.
Neutron irradiations were perR)rmed in JRR-2 and JMTR. The irradiation temperature and fluence ranged from 290 to 410°C, and from 0.88 to 4.2 × 10 > n / m (E > I MeV), respectively. Thc elastic plastic fracture toughness (Jm) was evaluated in principle following the multi-specimen method defined in the ASTM Test Method for J a , a measure of Fracture Toughness (E 813-87). DCT specimens were precracked by fatigue with a frequency of 10 Hz at room temperature such that the ratio of crack length to specimen width ( a / w ) was about 0.6. The final stress intensity factor during the precracking pro. cedure was kept to be less than 25 MPa~/m-. These specimens were then side grooved on each side to 12.5% of the thickness. The nmlti-specimen method was adopted because this method has a smaller error compared to the single specimen unloading compliance method in calculating the crack length, especially for small crack extensions. JIc tests were conducted at room temperature (20°C), which corresponded to the upper shelf region, leading to a completely ductile fracture for both unirradiated and irradiated CVN impact specimens. The tearing modulus was calculated using the equation T ma t =
(4)
(E/o'~)(dJ/da).
2.2.1. Effect o f neutron irradiation on J - R curve Figure 1 shows the J - R curves of base metal and welded joints before and after irradiation at about 400°C to a fluence of about 1×1023 n / m 2 ( E > l MeV). Before irradiation, the heat-affected zone (HAZ) had the lowest values of fracture toughness (Jlc) and tearing modulus (Tm~t). However, the decrease in both of these values caused by neutron irradi--
Table 1 Chemical composition of 2¼Cr-lMo steel C
A387 Gr.22 CI.2 (Normalized NT-1 0.15 NT-2 0.13 WELD (NT-1) 0.10
Si
Mn
and tempered) 0.05 0.55 0.10 0.55 0.20 0.66
P
S
Ni
Cr
Cu
Mo
As
Sn
0.008 0.005 0.005
0.0l 0.005 0.005
0.11 0.04 0.02
2.33 2.44 2.26
0.07 0.03 0.05
0.90 1.05 1.00
0.006 0.002 0.003
0.008 0.001 0.001
Heat treatment NT-1. 900/930°C - 6.5 h, water cooled, 670/690°C - 7 h, air cooled, PWHT (680/710°C - 20 h) NT-2, 920°C - 6.5 h, water cooled, 600°C - 6.5 h, air cooled, PWHT (685/695°C - 20 h)
Y. Muto et al. / High temperature structural materials for HTGR ation was the smallest for HAZ. As a result, the fracture toughness and tearing modulus after irradiation were almost the same for the base metal, weld metal and HAZ, and had an average value of 240 k J / m 2 for J[c (Jiep = J value at the crack initiation by electrical potential method = 160 k J / m 2) and 130 for Treat • It is known that the fracture toughness of pressure vessel steels decreases with increasing temperature [5,6]. Using the fracture toughness value 160 k J / m 2 and a fracture appearance transition temperature of - 2 5 ° C measured for the irradiated base metal, the fracture toughness at 400°C is estimated to be = 70 k J / m 2, on the basis of the method proposed in ref. [5].
2.2.2. Irradiation embrittlement o f thermally aged material Since 2 1 / 4 C r - l M o steel is to be exposed to a temperature of around 400°C for 105 h during the service life of the HTTR, it is of importance to evaluate the toughness degradation caused by thermal aging itself. A possible degradation would be age softening and temper embrittlement. The latter occurs in the temperature range from 375 to 550°C due to intergranular segregation of impurity atoms such as phosphorus and tin. Therefore, the combined effects of thermal aging and neutron irradiation embrittlement have to be examined for structural integrity. For thermally aged material, neutron irradiation caused only a slight decrease in the J - R curve level as compared with nontreated base plate material. To clarify this observation, Fig. 2 is shown to clearly illustrate the shift of Jlc and Tmat values with thermal aging and neutron irradiation. The decrease in both J[c and Tmat of the thermally aged material is much less than the unaged material, resulting in similar values for Jlc and 1000
800
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1.5
2
Crack extension Aa(mm) Fig. 1. J - R
curves for 2 ¼Cr-lMo steel before and after irradiation.
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400
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400
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Tearing modulus Fig. 2. Changes in J[c and the tearing modulus caused by thermal aging and neutron irradiation. rma t after irradiation in both materials. This different sensitivity to neutron irradiation is considered to reflect combined effects of thermal aging and neutron irradiation peculiar to elevated temperature irradiation. This point will be discussed in the next section.
2.2.3. Effect of irradiation temperature on the mechanical properties in the upper shelf region As was stated in the previous section, an interaction between thermal aging and neutron irradiation embrittlement may exist. After thermal aging at temperatures above 400°C, normalized and tempered 2 1 / 4 C r - l M o steel is subject to mechanical property changes depending on exposure time. For example, the DBTT (ductile to brittle transition temperature) begins to increase after aging at 400°C for 3 × 104 h due to phosphorus segregation to grain boundaries [7]. The effect of irradiation temperature on the mechanical properties in the upper shelf region will be discussed next. J - R curves for the materials irradiated at 290, 350 and 400°C to a fluence of about 1 to 3 x 1023 n / m 2 were measured: no change was caused by the irradiation at 350°C, whereas the curves shifted to the lower region for the other two irradiations. Figure 3 shows the percentage decrease of J]c, Jiep, Tmat and USE (Upper Shelf Energy) due to irradiation as a function of irradiation temperature, also including the above decrease due to thermal aging at 450°C for 1 × 10 4 h. When increasing irradiation temperature up to 350°C, the decrease in these fracture toughness parameters becomes less with an apparent minimum at around 350°C, and then increases again at 400°C. The decrease observed up to 350°C is coincident with a decrease in radiation hardening with increasing temperature.
Y. Muto et al. / High temperature structural materials for HTGR
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Exposure temperature (°C)
500
Thermally ag%d ~,. I material w
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Irradiation temperature (°C)
Fig. 3. Irradiation temperature dependence of the decrease of fracture toughness (Jiep: JIc determined by electropotential method, USE: Upper Shelf Energy).
Figure 4 shows the 0.2% offset stress as a function of i r r a d i a t i o n t e m p e r a t u r e . T h e indicated points are the average values of several test trials. T h e r a d i a t i o n h a r d e n i n g b e c o m e s very small at t e m p e r a t u r e s h i g h e r t h a n 350°C with a n e u t r o n fluence of ~ 1023 n / m 2. Nevertheless, some o t h e r factor s e e m e d to b e g i n to govern the d e c r e a s e of f r a c t u r e t o u g h n e s s at 400°C irradiation. J u d g i n g from t h e fact t h a t the d e c r e a s e of f r a c t u r e t o u g h n e s s values in t h e thermally aged m a t e rial (450°C, 1 x 104 h) by i r r a d i a t i o n was smaller t h a n the n o n t r e a t e d one, a n d t h a t t h e t h e r m a l aging at 450°C itself m a d e t h e t o u g h n e s s lower, t o u g h n e s s d e g r a d a t i o n caused by 400°C i r r a d i a t i o n is t h o u g h t to b e g o v e r n e d by a n i r r a d i a t i o n - e n h a n c e d thermally activ a t e d process.
Fig. 4. Irradiation temperature dependence of the increase in 0.2% offset stress caused by irradiation.
3. High-temperature alloys for structural applications 3.1. Development o f Hastelloy X R and its qualification test 3.1.1. Development o f Hastelloy X R R e f e r r i n g to a series of screening tests for c o m m e r cial a n d e x p e r i m e n t a l h e a t - r e s i s t a n t super-alloys, a nickel-base solid-solution s t r e n g t h e n i n g Hastelloy X was selected for application to the very h i g h - t e m p e r a ture c o m p o n e n t s in t h e H T T R . Hastelloy X, however, was j u d g e d to have n o sufficient compatibility with the primary coolant h e l i u m for long t e r m service. T h e r e fore the following modifications were m a d e to improve
Table 2 Chemical compositions of Hastelloy XR and Hastelloy X (mass%)
Hastelloy XR max rain Hastelloy X max min
Hastelloy XR max min Hastelloy X max min
C
Si
Mn
P
S
Cr
Co
Mo
W
0.15 0.05
0.50 0.25
1.00 0.75
0.040 -
0.030 -
23.0 20.5
2.50
10.0 8.0
1.00 0.20
0.15 0.05
1.00 -
1.00
0.040 -
0.030 -
23.0 20.5
2.50 0.50
10.0 8.0
1.00 0.2(}
Fe
Ni
B
AI
Ti
Cu
20.0 17.0
Bal. Bal.
0.010 -
0.05 -
0.03 -
0.50 -
20.0 17.0
Bal. Bal.
0.010
0.50 -
0.15 -
0.50 -
Y. Muto et al. / High temperature structural materials for HTGR long term performance in the service environment at elevated temperatures. (1) Optimizing Mn and Si contents to form a stable and adherent oxidation film [8], (2) Lowering A1 and Ti contents to suppress internal oxidation and intergranular attack [8], (3) Lowering Co content to minimize radioactive contamination in the primary system [8], and (4) Optimizing B content to improve the creep properties [9]. The specification of the improved version of Hastelloy X which is called the nuclear grade Hastelloy XR is shown in Table 2, with a comparison to that of Hastelloy X[10]. Figure 5 shows results of long-term corrosion tests in JAERI-type B helium environment (20 Pa H2, 10 Pa CO, 0.1 Pa H 2 0 , 0.2 Pa CO 2 and 0.5 Pa CH4, hereafter the helium environment) under severe thermal cyclings, where superiority of Hastelloy XR is demonstrated relative to Hastelloy X [11].
3.1.2. ComprehensiL,e qualification test for Hastelloy XR Comprehensive qualification tests such as creep, fatigue, corrosion and other fracture-relevant properties on Hastelloy XR have been carried out in order to accumulate the test data for structural design and safety evaluation. Figure 6 shows the status of on-going creep tests with a complete record of deformation characteristics in the helium environment at temperatures ranging from 800 to 1050°C and some of the creep tests are being continued exceeding 4.7 × 104 h. From this figure, trends in stress dependence and data scattering of the creep rupture strength are judged to be quite similar at 1000°C to those at lower temperatures. Therefore, Hastelloy XR can be concluded to be
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stable at 1000°C or below [12]. As for fatigue tests, the effects of strain rate [13], hold time [13], aging [14] and test temperatures[15] on fatigue properties have been studied in the helium environment. Corrosion tests revealed that Hastelloy XR still showed stable corrosion characteristics in maintaining protective oxide integrity without significant spallation or intergranular attack even after exposure to a helium environment at 900°C for 3 × 104 h under severe thermal cycling as shown previously.
3.2. Recent adL,ancements on high-temperature alloys 3.2.1. Development o f a filler metal for Hastelloy XR weldment The strengths, especially creep strength, of the weldments with conventional filler metals at high temperatures are lower than those of Hastelloy XR base metal [16]. Therefore a series of exploratory studies was performed to develop a filler metal for Hastelloy XR weldments with the aim of possessing excellent weldability and high temperature strength properties simultaneously. In the first step, an alloy design method was organized on the basis of multiple regression analysis by using the data available in open literature. Validity of the proposed method and its usefulness for developing filler metals were confirmed experimentally through weldability, tensile and creep tests [17]. In the second step, to improve the susceptibility of cracking in the weldments without degrading the high temperature strength properties, optimization of B and C contents of the filler metals within the specification range of Hastelloy XR was attempted. As shown in Fig. 7, the susceptibility of cracking in the weldments
310
}i.. Mute et a L / High temperature structural materials~or HTGR
i
z 5_~
;J / /
//0.0025
could be substantially lowered with an adjustment of chemical composition of the filler metals. And the creep strengths of the tungsten-inert-gas (TIG) weldments with the improved filler metals are higher than those of Hastelloy XR base metals as shown in Fig. 8. Therefore, it is concluded that filler metals developed during the course of this study are applicable to Hastelloy X R welded-structures for the H T T R since the weldments with the newly developed filler metals showed an excellent performance.
3.2.2. Det,elopment of N i - C r - W superalloy for future HTGR application As to the development of N i - C r - W superalloy for future H T G R applications with coolant outlet temperatures of around 1000°C, the program had set the following targets: (1) Creep rupture strength over 7.8 MPa at 1000°C for 1 x 1 0 s hlife, (2) No substantial susceptibility to intergranular attack or internal oxidation after exposure to the helium environment at 1000°C for 1 x 105 h, (3) Formability to 32-ram-outer diameter, 5-mm-thick, 7-m-long seamless tubes, and (4) Sufficient capacity for the secondary cold working and welding. In the first step of the program, the basic chemical composition, i.e., the C r / W ratio of the alloy, was examined. In the N i - C r - W alloy system, the strengthening due to the W-rich bcc Gz-W phase is most effective when it precipitates in a manner covering the grain boundaries. Hence creep rupture strength, tensile property, hot workability and corrosion resistance in the helium environment were examined on five N i - C r - W alloys, from 28 mass% Cr/15.5 mass% W to 12 mass% C r / 2 7 . 5 mass% W, in which the W rich bee
' ....
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Fig. 7. Effect of B and C contents of the filler metals on the susceptibility of cracking in TIG weldments through side bending test.
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•
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t0 2
5
Creep
t0 3
2
life
rupture
5
t0 4
(h}
Fig. 8. Relation between applied stress and time to rupture for Hastelloy XR weldments at 900°C in air. Solid and broken lines are mean and design values of creep rupture strength for Hastelloy XR, respectively (WM: Weld metal, WJ: Welded joint).
a2-W phase precipitates in such a manner. And the basic chemical composition of the alloy was determined to be N i / 1 8 to 19 mass% C r / 2 0 to 22 mass% W mainly from the viewpoints of creep rupture strength and corrosion resistance in a helium environment [18]. In the second step, effects of the further alloying elements, i.e., C, Nb, Fe, Mn, Si, Ti, B and Y were investigated systematically on creep rupture strength, tensile properties before and after aging, hot workability and corrosion resistance in a helium environment. A series of experiments has revealed that the addition of Nb, Fe a n d / o r enrichment of C more than 0.03
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2
3
oxygen pickup (mil/cm 2)
Fig. 9. Relation between experimental and predicted oxygen pickup (OP) after exposure to HTGR helium environment at 1000°C for 1000 h. Predicted OP=0.8269+0.5684Mn1.9013Si-74.8233B-53.8198¥+3.3176Ti where OP is in mg/cm 2, and Mn, Si, B, Y and Ti are in mass%.
311
Y. Muto etal. / High temperature structural materials for HTGR Creep rupture strength
,ooo i
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(tO00%, 30MPo )
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Ni-Cr-W superolloy/ !1' , \
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Corrosion resistance
RT tensileductilityoffer aging
(HTGR helium, tO00°C,lO00h)
(Elong.ofter aging at 800 or 900"C for lO00h )
Fig. 10. Comparison of newly developed superalloy (predicted values) with the existing high-temperature alloys. mass% exhibits no beneficial influence and that the addition of B and Y in either singular or simultaneous forms is remarkably efficient in improving either creep rupture strength, tensile ductility, corrosion resistance or hot workability. As for the effect of Mn, Si and Ti, they have both beneficial and detrimental effects. Figure 9 is an example of multiple regression analysis results that shows the relation between experimental and predicted oxygen pickup after exposure to a helium environment at 1000°C for 1000 h [19]. Based on the experimental results, it has been judged that Ni/18 to 19 mass% Cr/20 to 22 mass% W/0.03 mass% C/0.08 mass% Ti/0.02 to 0.05 mass% Zr/0.002 to 0.007 mass% Y/0.0035 to 0.006 mass% B is the optimum chemical composition. And the newly proposed N i - C r - W superalloy can be expected to in-
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20 30 40 Time (h) Fig. 12. Comparison between experimental data and calculated creep curves.
dicate well-balanced properties and superior performance to the existing high-temperature alloys as illustrated in Fig. 10.
4. Structural test for the high temperature design code In order to develop an advanced life prediction method at the high temperature exceeding around 800°C, developing both a constitutive equation and an improved creep-fatigue damage evaluation method is required. Then, efforts are focussed on it. 4.1. Constitutive equation
Constant stress-creep tests were conducted at 900°C for Hastelloy XR2 which is a modified version of HastelloyXR. A uni-axial creep constitutive equation of the Garofalo type was made which had three parameters such as Et: asymptotic primary strain, r: recipro-
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(7 (MPo) Fig. 11. Results of stress dip tests for Hastelloy XR2.
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Fig. 13. Prediction by the time fraction rule.
312
Y. Muto et al. / High temperature structural materials for HTGR r
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° 0
Fig. 14. Prediction by the ductility exhaustion rule. cal of time constant and groin: minimum creep strain rate [20]. The values of the parameters were determined experimentally. The parameters r and Emin were fitted very well, though a significant data-scattering was observed in the parameter Et. Therefore, it can be said that this type of creep equation must be one of the most adequate ones for Hastelloy XR2. The equation was applied to predict a creep behavior under a variable stress condition using the strain hardening law. The result was, however, insufficient and suggested the effect of an internal back stress. Then, a new constitutive equation incorporating the internal back stress was made where a creep strain rate equation was assumed to be the well-known power law type [21], and an evolution equation of the internal back stress was based on a creep-recovery model [22]. The parameters were determined by stress dip tests whose results are shown in Fig. 11. The new constitutive equation is given as follows [23]. - ~ b ( ~ ) ]~('~',
(5)
6% = B ( ~ ) ~ - D(o-)o-~ °s3
(6)
= A(~)[~
O: 900
v
0
/j
100
!
l
1
I
I
i
200
300
400
500
600
Time (h) Fig. 16. Biaxial creep strain behavior.
where A 0 r ) , B(~r) and D((r) are material constants, (rb. 0 = an initial value of the internal back stress. The creep behaviors were calculated by both the
E
o
10 3
Nf,r: Nf,p Y o 95o'cx.r.
q3 t 0 2 I
t0 2
.85o'cx,,T.
950"C XRCH • 850°CXR, CH [] 950*(: XR TCH • 850°CXRIITCH
¢J
I
I
Jllll~
I
I
J J;ll[J
103
t0 4
Real cycles to failure Nf,r Fig. 15. Prediction by the damage rate equation.
Fig. 17. Cyclic thermal loading of cylindrical specimen.
Y. Muto et al. / High temperature structural materials for HTGR
new equation and the Garofalo equation for variable stress loads. It was ascertained that the new constitutive equation could predict the creep behavior under the variable stress loads more exactly than the Garofalo equation as shown in Fig. 12. A stress-relaxation behavior is, however, frequently observed in thermal loading cycles and a creep damage during the relaxation becomes important. Therefore, further work is needed to examine the validity of the new creep constitutive equation for the relaxation. 4.2. Creep-fatigue tests
The low cycle fatigue tests with and without a dwell period were conducted at 850°C and 950°C under vacuum conditions of 10 3 Pa. The strain rate in loading and unloading is changed from 0.01%/s to 0 . 5 % / s at 850°C, but to l % / s at 950°C. The dwell period in a cycle ranges from 150 s to 1800 s at tension a n d / o r compression sides: this time duration is enough to achieve the low level of residual stress rate where the cycles to failure won't be changed. Cycles to failures were predicted by the time fraction rule [24], a ductility exhaustion rule [25] and a damage rate equation [26], whose results are shown in Figs. 13, 14 and 15, respectively. In the time fraction rule the predicted cycles to failure were larger than the real ones in TH and CH
313
cycles: TH is a cycle with a dwell period in tension and CH in compression. Consequently a failure will occur at the creep-fatigue damage less than unity in these cycles. The ductility exhaustion rule could take the effect of strain rate into account. The method is more likely to reflect the mechanism of creep damage rather than the time fraction rule though the accuracy of preditions by both methods are the same. It is, however, still suggested from the results that the damage is too much counted in TCH cycle: TCH is a symmetrical cycle with a dwell period in tension and compression sides. A concept of the damage-rate equation is basically the growth of cavities, described as a function of inelastic strain range. The damage or the cavity growth in the CH cycle will need further discussion. 4.3. Biaxial test
Biaxial stress-rupture tests were conducted for the condition of both the combined tension and torsion loads and the combined compression and torsion loads. It was ascertained that a uniaxial tensile stress-rupture test results in the shortest rupture life for the same equivalent stress of constant loading conditons [12]. Then, tests were succeeded under constant equivalent stress conditions. Thin-walled cylindrical test specimen of 12 mm gauge length, 22 mm outer diameter and 2 mm wall thickness were used.
Fig. 18. Cyclic bending of curved tube.
314
Y. Muto et aL / High temperature structural materials fi)r HT(;R
Creep strain behaviors are shown in Fig. 16 by a parameter of angle in the stress space. The creep strain rates in the primary creep region is retarded as the angle decreased. The data of creep strain exceeding several percents are not reliable because those were accelerated by buckling. Buckling due to the eccentricity of the loading axis and bulging type bucklings were observed in the case of torsion and compression, respectively. Biaxial creep-fatigue tests with constant torsion and cyclic axial load were conducted at 900°C. The experimental creep-fatigue life coincided with the analyrical one predicted by the time fraction rule. 4.4. Component test
a thick wall will take thermally cycled loading and a heat transfer tube will take mechanically cycled loading. Photos of the former test section and the latter test apparatus are shown in Figs. 17 and 18, respectively. In curved tube-cyclic bending tests at 900°C, cracks were observed to initiate at the bottom of the curved corner and run to the circumferential direction in the case of in-plane bending tests, while they initiated at the side and propagated to the direction of 45 ° to the tube axis in the case of out-of-plane bending test. The results coincide with those by the stress analyses as shown in Fig. 19. 5. Conclusions
The limits of thermal loading to the progressive time-dependent damage of high temperature components in the intermediate heat exchanger system have heen ~tndied in mnlti-n~zial ,~tre,:~ ~t~te~" a cylinder with
(1) For 2 l / 4 C r - l M o steel the fracture toughness at 400°C, the service temperature of the pressure vessel, was estimated to be = 70 k J / m z after irradiation, ten times as high as H T T R life fluence.
/
I Maximumshear stress {n-plane bending \ (Out-of-planebending~/~ (/ (Experiment) Out-of-planebending ><~ ,/VI '~ ..... (Experiment) ~,.~. ~ ,) Loaa,ng s,ae / [
" [ \-.. I Maximumlongitudinalstress / ~J (In-planebending(Analysis)) Fixed side Fig. 19. Appearance of failure tubes.
Y. Muto et al. / High temperature structural materials )br HTGR (2) T h e d e c r e a s e in the fracture t o u g h n e s s of thermally aged material caused by irradiation is m u c h less t h a n that of the as-received base plate material. (3) T h e radiation h a r d e n i n g b e c a m e very small at temp e r a t u r e s higher than 350°C with a n e u t r o n fluence of a b o u t 3 × 1023 n / m 2 ( E > 1 MeV). (4) Hastelloy X R base metals as well as filler metals for the weldings have b e e n developed and characterized to apply to the high t e m p e r a t u r e compon e n t s of the H T T R . (5) O p t i m u m chemical composition of newly developed N i - C r - W superalloy, which can be expected to indicate well-balanced properties, has b e e n proposed. (6) T h e constitutive e q u a t i o n incorporating the internal back stress could r e p r e s e n t the creep strain b e h a v i o r of Hastelloy X R at variable stress conditions. (7) Applicabilities of the time fraction rule, the ductility exhaustion rule and the d a m a g e rate e q u a t i o n to Hastelloy X R were e x a m i n e d b a s e d on experim e n t a l data. (8) Biaxial c r e e p a n d creep-fatigue data were o b t a i n e d for Hastelloy XR.
References [1] [2] [3] [4]
S. Saito et al., Nucl. Engrg. Des. 132 (1991) 85. T. Kodaira et al., JAERI-M 85-170 (1985). T. Oku et al., Nucl. Engrg. Des. 119 (1990) 177. Y. Nishiyama et al., ASTM STP 1125, R.E. Stroller et al., Eds. (1992) pp. 1287-1303. [5] W.A. Logsdon and J.A. Begley, Eng. Fracture Mechs. 9 (1977) 461.
315
[6] T. Iwadate et al., ASTM STP 631, J.M. Barsom, Eds. (1977) pp. 493-506. [7] M. Suzuki et al., ASTM STP 956, F.A. Garner et al., Eds. (1987) pp. 98-110. [8] M. Shindo and T. Kondo, Proc. Int. Conf. on Gas-Cooled Reactor Today, Bristol, UK, 1982 (British Nuclear Energy Society) 2, pp.179-184. [9] Y. Kurata et al., Proc. Int. Conf. on Creep, Tokyo, Japan, 1986 (The Japan Society of Mechanical Engineers), pp. 97-102. [10] K. Hada et al., JAERI-M 911-148 (1990). [11] T. Tsukada et al., Proc. Workshop of Advanced Materials and Protective Coatings, Tokyo, Japan, 1992 (Elsevier Science Publishers) pp. 233-242. [12] K.Hada et al., Nucl. Engrg. Des. 132 (1991) 1-11. [13] H. Tsuji and T. Kondo, J. Nucl. Mater. 151t (1987) 259265. [14] H. Tsuji and H. Nakajima, J. Nucl. Mater. 158 (1988) 267-278. [15] H. Tsuji and H. Nakajima, J. Nucl. Mater. 151 (1987) 1-9. [16] T. Udoguchi and T. Nakanishi, Int. J. Press. Vessel and Piping 9 (19811 107-123. [17] K. Watanabe et al., J. Nucl. Mater. 185 (19911 8-18. [18] Ad Hoc Committee on Advanced Superalloys, Technical Expert Committee on HTGR, Japan Atomic Energy Research Institute, JAERI-M 88-270 (1989). [19] H. Tsuji et al., JAERI-M 91-136 (19911. [20] F.Garofalo, Fundamentals of Creep and Creep-Rupture in Metals (Macmillan, New York, 1965). [21] F.H. Norton, Creep of Steel at High Temperatures (McGraw Hill, New York, 1929). [22] R. Lagneborg, Metal Science J. 3 (1969) 161-168. [23] Y. Kaji et al., The 5th Int. Conf. on Creep of Materials (1992). [24] Cases of ASME Boiler and Pressure Vessel Code, Case N-47. [25] R.H. Priest and E.G. Ellison, Material Science and Engrg. 49 (19811 7-17. [26] S. Majumdar and P.S. Maiya, J. Engrg. Materials & Technol. 102 (1980) 159-167.
305
Development of high temperature structural materials for the HTGR Y. Muto, H. N a k a j i m a a n d M. Eto Tokai Research Establishment, Japan Atomic Energy Research Institute, 2-4 Shirane, Shirakata, Tokai-mura, lbaraki-ken, 319-11, Japan Received 24 May 1993
In accordance with the HTGR program in Japan, a series of R&D for high temperature structural materials in particular with respect to the HTTR design code has been performed in JAERI for more than 20 years. This paper introduces R&D results of the pressure retaining low alloy steel 2 1/4Cr-lMo and the high temperature structural alloys Hastelloy XR and Ni-Cr-W superalloy for the design code together with some fruits of recent studies.
1. Introduction Studies on metallic materials for high temperature gas-cooled reactor (HTGR) applications have been made at the Japan Atomic Energy Research Institute(JAERI) for more than 20 years. A high temperature design code for the High Temperature Engineering Test Reaetor(HTTR)[1] has been already organized based on the data obtained by a series of material testings and structural tests. The items of research activities carried out are classified into three categories, i.e., those for the pressure retaining low alloy steels, for the characterization of the high-temperature alloys and for their structural application at very high temperature. This paper presents an overview of the JAERI's program during the last decade and emphasizes recent substantial advancement in each category.
2. Development of alloys for the pressure boundary structural components for the HTrR Among various components in the nuclear reactor, the pressure vessel is the one which requires a very high reliability. During normal operation of the HTTR the temperature of the vessel is supposed to be 400°C. However, there would be a possibility that the temperature would rise to 440°C at an abnormal transient state so that the design temperature was set at this temperature. It was decided that 2 1 / 4 C r - l M o steel
would be employed for the pressure vessel of the HTTR in consideration of the higher service temperature comparing with the light water reactor [2]. R & D of the steel have been carried out focussing on the aging behavior in relation to the chemical composition, rolling and forging process, welding method, etc.. To characterize the material, tensile and fatigue strengths, fracture toughness, Charpy impact energy etc. have been measured for the irradiated and thermally aged specimens. Here, the results of characterization tests are summarized together with some fruits of recent studies. 2.1. Application of 2 1 / 4 Cr-lMo steel to the HTTR Comparison between the normalized and tempered (NT) material and the annealed (AN) one was made from various aspects. As a result, NT was selected on the basis of the following facts. (1) Both strength and toughness increase due to the NT treatment which also rises the design stress up to around 450°C. (2) Higher toughness is maintained at the lowest service temperature. (3) NT is more resistant to the radiation and thermal aging embrittlements. To specify the NT material for the pressure vessel the following procedures were taken on the basis of the data on strength, ductility and toughness of the thick plate [3]. (1) With regard to the thickness ( ~ 160 mm) of the plate, a quenched and tempered (QT) treatment equiv-
0 0 2 9 - 5 4 9 3 / 9 3 / $ 0 6 . 0 0 © 1993 - Elsevier Science Publishers B.V. All rights reserved
Y Muto et al. / High temperature structural materials for HTGR
306
alent to NT was employed. The optimum post weld heat treatment (PWHT) parameter P was determined as
P = T(log t + 20) × 10 -3 = 2 0 . 5 -
20- 75,
(1)
where T and t are temperature and time in °C and hr, respectively. The effect of the normalizing treatment can be disregarded here because the NT temperature is set a little bit lower than usual so that the virtually PWHT is only an influential treatment. (2) For the reduction of thermal aging and radiation embrittlements, the requirements on the chemical composition were derived empirically, i.e., J = (Si + M n )( P + Sn) × 104 < 100,
(2)
10P + 5Sb + 4Sn + As =
100
< 10.
(3)
Here, Si, Mn, P, etc. are expressed in mass%. These requirements indicate that not only impurity elements such as Si and P but also alloying elements such as Mn and Sn are important in the determination of the embrittlement parameter.
2.2. Effect o f neutron irradiation embrittlement on the thermal aging embrittlement The neutron irradiation embrittlement of the NT material in the upper shelf region as well as the effect of neutron irradiation on thermal aging embrittlement was investigated in terms of elastic-plastic fracture toughness. The material used is shown in Table 1 [4]. Irradiated specimens were 4 mm diameter by 26 mm gauge length tensile specimens, standard size Charpy V-notch (CVN) specimens, and disk-shaped compact tension (DCT) specimens with three different thickness values.
Neutron irradiations were perR)rmed in JRR-2 and JMTR. The irradiation temperature and fluence ranged from 290 to 410°C, and from 0.88 to 4.2 × 10 > n / m (E > I MeV), respectively. Thc elastic plastic fracture toughness (Jm) was evaluated in principle following the multi-specimen method defined in the ASTM Test Method for J a , a measure of Fracture Toughness (E 813-87). DCT specimens were precracked by fatigue with a frequency of 10 Hz at room temperature such that the ratio of crack length to specimen width ( a / w ) was about 0.6. The final stress intensity factor during the precracking pro. cedure was kept to be less than 25 MPa~/m-. These specimens were then side grooved on each side to 12.5% of the thickness. The nmlti-specimen method was adopted because this method has a smaller error compared to the single specimen unloading compliance method in calculating the crack length, especially for small crack extensions. JIc tests were conducted at room temperature (20°C), which corresponded to the upper shelf region, leading to a completely ductile fracture for both unirradiated and irradiated CVN impact specimens. The tearing modulus was calculated using the equation T ma t =
(4)
(E/o'~)(dJ/da).
2.2.1. Effect o f neutron irradiation on J - R curve Figure 1 shows the J - R curves of base metal and welded joints before and after irradiation at about 400°C to a fluence of about 1×1023 n / m 2 ( E > l MeV). Before irradiation, the heat-affected zone (HAZ) had the lowest values of fracture toughness (Jlc) and tearing modulus (Tm~t). However, the decrease in both of these values caused by neutron irradi--
Table 1 Chemical composition of 2¼Cr-lMo steel C
A387 Gr.22 CI.2 (Normalized NT-1 0.15 NT-2 0.13 WELD (NT-1) 0.10
Si
Mn
and tempered) 0.05 0.55 0.10 0.55 0.20 0.66
P
S
Ni
Cr
Cu
Mo
As
Sn
0.008 0.005 0.005
0.0l 0.005 0.005
0.11 0.04 0.02
2.33 2.44 2.26
0.07 0.03 0.05
0.90 1.05 1.00
0.006 0.002 0.003
0.008 0.001 0.001
Heat treatment NT-1. 900/930°C - 6.5 h, water cooled, 670/690°C - 7 h, air cooled, PWHT (680/710°C - 20 h) NT-2, 920°C - 6.5 h, water cooled, 600°C - 6.5 h, air cooled, PWHT (685/695°C - 20 h)
Y. Muto et al. / High temperature structural materials for HTGR ation was the smallest for HAZ. As a result, the fracture toughness and tearing modulus after irradiation were almost the same for the base metal, weld metal and HAZ, and had an average value of 240 k J / m 2 for J[c (Jiep = J value at the crack initiation by electrical potential method = 160 k J / m 2) and 130 for Treat • It is known that the fracture toughness of pressure vessel steels decreases with increasing temperature [5,6]. Using the fracture toughness value 160 k J / m 2 and a fracture appearance transition temperature of - 2 5 ° C measured for the irradiated base metal, the fracture toughness at 400°C is estimated to be = 70 k J / m 2, on the basis of the method proposed in ref. [5].
2.2.2. Irradiation embrittlement o f thermally aged material Since 2 1 / 4 C r - l M o steel is to be exposed to a temperature of around 400°C for 105 h during the service life of the HTTR, it is of importance to evaluate the toughness degradation caused by thermal aging itself. A possible degradation would be age softening and temper embrittlement. The latter occurs in the temperature range from 375 to 550°C due to intergranular segregation of impurity atoms such as phosphorus and tin. Therefore, the combined effects of thermal aging and neutron irradiation embrittlement have to be examined for structural integrity. For thermally aged material, neutron irradiation caused only a slight decrease in the J - R curve level as compared with nontreated base plate material. To clarify this observation, Fig. 2 is shown to clearly illustrate the shift of Jlc and Tmat values with thermal aging and neutron irradiation. The decrease in both J[c and Tmat of the thermally aged material is much less than the unaged material, resulting in similar values for Jlc and 1000
800
/
/
m
e
t
a
Unirradiateddataband l
e"E 600 eA
Q&&
v400
"-~
--
200
0
~ I 0.5
&&A Testtemp. : R.T. lrrad,temp. : 393~403"C Fluence : 0.74~1.1 X 1023n/m~ OBase AHAZ mWdd iOpensymbols : lunirr. 1
1.5
2
Crack extension Aa(mm) Fig. 1. J - R
curves for 2 ¼Cr-lMo steel before and after irradiation.
500
400
i
307 i
i
i
r . 1Aged + lrrad-) k~* 2 Irrad. /
~E" 300 v
-.j
* 1~.~"/
200
Initial state
--~Aging
(450"C × lO000h)
*~2
Unirr. ~ Irr. (400"(2, 3 X 1023n/m 2) 100
, 06
I 1 00
,
I 200
0 Initialstate (as received [] Pre-inad.aged , •lhradia,~d l , 300
400
500
Tearing modulus Fig. 2. Changes in J[c and the tearing modulus caused by thermal aging and neutron irradiation. rma t after irradiation in both materials. This different sensitivity to neutron irradiation is considered to reflect combined effects of thermal aging and neutron irradiation peculiar to elevated temperature irradiation. This point will be discussed in the next section.
2.2.3. Effect of irradiation temperature on the mechanical properties in the upper shelf region As was stated in the previous section, an interaction between thermal aging and neutron irradiation embrittlement may exist. After thermal aging at temperatures above 400°C, normalized and tempered 2 1 / 4 C r - l M o steel is subject to mechanical property changes depending on exposure time. For example, the DBTT (ductile to brittle transition temperature) begins to increase after aging at 400°C for 3 × 104 h due to phosphorus segregation to grain boundaries [7]. The effect of irradiation temperature on the mechanical properties in the upper shelf region will be discussed next. J - R curves for the materials irradiated at 290, 350 and 400°C to a fluence of about 1 to 3 x 1023 n / m 2 were measured: no change was caused by the irradiation at 350°C, whereas the curves shifted to the lower region for the other two irradiations. Figure 3 shows the percentage decrease of J]c, Jiep, Tmat and USE (Upper Shelf Energy) due to irradiation as a function of irradiation temperature, also including the above decrease due to thermal aging at 450°C for 1 × 10 4 h. When increasing irradiation temperature up to 350°C, the decrease in these fracture toughness parameters becomes less with an apparent minimum at around 350°C, and then increases again at 400°C. The decrease observed up to 350°C is coincident with a decrease in radiation hardening with increasing temperature.
Y. Muto et al. / High temperature structural materials for HTGR
308 100
V
/~,
80 ~ ¢o
60
,
~
3
~
•
o
•
250
.... .-. . . . . . . . .
~4o
i
Test temp.=R,T. Fluence=l ~ 3 x l~'3rt/m2~> lMe\)
~30
Thermallyaged rnateriall
el
o ~20
,
•
F
~1o
~
~
i
II
J
300
350
400
i-
A
I$"
• o
---r . . . . . .
>
./
OUSE
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'
Fluence=l ~ 3 X 1~ n/m (E IMeV) &'JIc Open symbols~without irradiation ! "Jiep (450"C X ]0000h aged) q i • 1Trnat
o
;
b
i__J 450
Exposure temperature (°C)
500
Thermally ag%d ~,. I material w
I 250
300
350
400
450
Irradiation temperature (°C)
Fig. 3. Irradiation temperature dependence of the decrease of fracture toughness (Jiep: JIc determined by electropotential method, USE: Upper Shelf Energy).
Figure 4 shows the 0.2% offset stress as a function of i r r a d i a t i o n t e m p e r a t u r e . T h e indicated points are the average values of several test trials. T h e r a d i a t i o n h a r d e n i n g b e c o m e s very small at t e m p e r a t u r e s h i g h e r t h a n 350°C with a n e u t r o n fluence of ~ 1023 n / m 2. Nevertheless, some o t h e r factor s e e m e d to b e g i n to govern the d e c r e a s e of f r a c t u r e t o u g h n e s s at 400°C irradiation. J u d g i n g from t h e fact t h a t the d e c r e a s e of f r a c t u r e t o u g h n e s s values in t h e thermally aged m a t e rial (450°C, 1 x 104 h) by i r r a d i a t i o n was smaller t h a n the n o n t r e a t e d one, a n d t h a t t h e t h e r m a l aging at 450°C itself m a d e t h e t o u g h n e s s lower, t o u g h n e s s d e g r a d a t i o n caused by 400°C i r r a d i a t i o n is t h o u g h t to b e g o v e r n e d by a n i r r a d i a t i o n - e n h a n c e d thermally activ a t e d process.
Fig. 4. Irradiation temperature dependence of the increase in 0.2% offset stress caused by irradiation.
3. High-temperature alloys for structural applications 3.1. Development o f Hastelloy X R and its qualification test 3.1.1. Development o f Hastelloy X R R e f e r r i n g to a series of screening tests for c o m m e r cial a n d e x p e r i m e n t a l h e a t - r e s i s t a n t super-alloys, a nickel-base solid-solution s t r e n g t h e n i n g Hastelloy X was selected for application to the very h i g h - t e m p e r a ture c o m p o n e n t s in t h e H T T R . Hastelloy X, however, was j u d g e d to have n o sufficient compatibility with the primary coolant h e l i u m for long t e r m service. T h e r e fore the following modifications were m a d e to improve
Table 2 Chemical compositions of Hastelloy XR and Hastelloy X (mass%)
Hastelloy XR max rain Hastelloy X max min
Hastelloy XR max min Hastelloy X max min
C
Si
Mn
P
S
Cr
Co
Mo
W
0.15 0.05
0.50 0.25
1.00 0.75
0.040 -
0.030 -
23.0 20.5
2.50
10.0 8.0
1.00 0.20
0.15 0.05
1.00 -
1.00
0.040 -
0.030 -
23.0 20.5
2.50 0.50
10.0 8.0
1.00 0.2(}
Fe
Ni
B
AI
Ti
Cu
20.0 17.0
Bal. Bal.
0.010 -
0.05 -
0.03 -
0.50 -
20.0 17.0
Bal. Bal.
0.010
0.50 -
0.15 -
0.50 -
Y. Muto et al. / High temperature structural materials for HTGR long term performance in the service environment at elevated temperatures. (1) Optimizing Mn and Si contents to form a stable and adherent oxidation film [8], (2) Lowering A1 and Ti contents to suppress internal oxidation and intergranular attack [8], (3) Lowering Co content to minimize radioactive contamination in the primary system [8], and (4) Optimizing B content to improve the creep properties [9]. The specification of the improved version of Hastelloy X which is called the nuclear grade Hastelloy XR is shown in Table 2, with a comparison to that of Hastelloy X[10]. Figure 5 shows results of long-term corrosion tests in JAERI-type B helium environment (20 Pa H2, 10 Pa CO, 0.1 Pa H 2 0 , 0.2 Pa CO 2 and 0.5 Pa CH4, hereafter the helium environment) under severe thermal cyclings, where superiority of Hastelloy XR is demonstrated relative to Hastelloy X [11].
3.1.2. ComprehensiL,e qualification test for Hastelloy XR Comprehensive qualification tests such as creep, fatigue, corrosion and other fracture-relevant properties on Hastelloy XR have been carried out in order to accumulate the test data for structural design and safety evaluation. Figure 6 shows the status of on-going creep tests with a complete record of deformation characteristics in the helium environment at temperatures ranging from 800 to 1050°C and some of the creep tests are being continued exceeding 4.7 × 104 h. From this figure, trends in stress dependence and data scattering of the creep rupture strength are judged to be quite similar at 1000°C to those at lower temperatures. Therefore, Hastelloy XR can be concluded to be
10 3
~
,
_~
,
~f,,~
I
tooo%/~
/ o /
~ to'
f
°
t
irt,,[
f
~o
g~
2 HTGR helium t0' t0;,
.......
'1'03
.......
1104
'
(h) Fig. 5. Comparison of environmental effect in Cr-depleted Time
zone depth between Hastelloy XR and Hastelloy X.
2
I
402 =E
I l I rill]
309
I I [ITll I
I
Hostelloy XR -HTGR helium
v
I flllll
I
i
I llllll
I
i
1 fllfl
l::gD O
N
~'I'~AO 0
5
•
~
Olo
0
. - ~ ~ =
oo
7~t~ ~."qK~~ . [] o •
< t0 t
800 900 950 1000 1050
• I
400
I IIIII1~
1o'
I
'
z~
IIII1111
I
402
~:~ A~"
I IIIIIll
L
1o3
•
o-,.o.-,..
IIiiiid
i
LIII
40"
to 5
Creep rupture life (h) Fig. 6. Creep rupture life for Hastelloy XR.
stable at 1000°C or below [12]. As for fatigue tests, the effects of strain rate [13], hold time [13], aging [14] and test temperatures[15] on fatigue properties have been studied in the helium environment. Corrosion tests revealed that Hastelloy XR still showed stable corrosion characteristics in maintaining protective oxide integrity without significant spallation or intergranular attack even after exposure to a helium environment at 900°C for 3 × 104 h under severe thermal cycling as shown previously.
3.2. Recent adL,ancements on high-temperature alloys 3.2.1. Development o f a filler metal for Hastelloy XR weldment The strengths, especially creep strength, of the weldments with conventional filler metals at high temperatures are lower than those of Hastelloy XR base metal [16]. Therefore a series of exploratory studies was performed to develop a filler metal for Hastelloy XR weldments with the aim of possessing excellent weldability and high temperature strength properties simultaneously. In the first step, an alloy design method was organized on the basis of multiple regression analysis by using the data available in open literature. Validity of the proposed method and its usefulness for developing filler metals were confirmed experimentally through weldability, tensile and creep tests [17]. In the second step, to improve the susceptibility of cracking in the weldments without degrading the high temperature strength properties, optimization of B and C contents of the filler metals within the specification range of Hastelloy XR was attempted. As shown in Fig. 7, the susceptibility of cracking in the weldments
310
}i.. Mute et a L / High temperature structural materials~or HTGR
i
z 5_~
;J / /
//0.0025
could be substantially lowered with an adjustment of chemical composition of the filler metals. And the creep strengths of the tungsten-inert-gas (TIG) weldments with the improved filler metals are higher than those of Hastelloy XR base metals as shown in Fig. 8. Therefore, it is concluded that filler metals developed during the course of this study are applicable to Hastelloy X R welded-structures for the H T T R since the weldments with the newly developed filler metals showed an excellent performance.
3.2.2. Det,elopment of N i - C r - W superalloy for future HTGR application As to the development of N i - C r - W superalloy for future H T G R applications with coolant outlet temperatures of around 1000°C, the program had set the following targets: (1) Creep rupture strength over 7.8 MPa at 1000°C for 1 x 1 0 s hlife, (2) No substantial susceptibility to intergranular attack or internal oxidation after exposure to the helium environment at 1000°C for 1 x 105 h, (3) Formability to 32-ram-outer diameter, 5-mm-thick, 7-m-long seamless tubes, and (4) Sufficient capacity for the secondary cold working and welding. In the first step of the program, the basic chemical composition, i.e., the C r / W ratio of the alloy, was examined. In the N i - C r - W alloy system, the strengthening due to the W-rich bcc Gz-W phase is most effective when it precipitates in a manner covering the grain boundaries. Hence creep rupture strength, tensile property, hot workability and corrosion resistance in the helium environment were examined on five N i - C r - W alloys, from 28 mass% Cr/15.5 mass% W to 12 mass% C r / 2 7 . 5 mass% W, in which the W rich bee
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Fig. 7. Effect of B and C contents of the filler metals on the susceptibility of cracking in TIG weldments through side bending test.
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Fig. 8. Relation between applied stress and time to rupture for Hastelloy XR weldments at 900°C in air. Solid and broken lines are mean and design values of creep rupture strength for Hastelloy XR, respectively (WM: Weld metal, WJ: Welded joint).
a2-W phase precipitates in such a manner. And the basic chemical composition of the alloy was determined to be N i / 1 8 to 19 mass% C r / 2 0 to 22 mass% W mainly from the viewpoints of creep rupture strength and corrosion resistance in a helium environment [18]. In the second step, effects of the further alloying elements, i.e., C, Nb, Fe, Mn, Si, Ti, B and Y were investigated systematically on creep rupture strength, tensile properties before and after aging, hot workability and corrosion resistance in a helium environment. A series of experiments has revealed that the addition of Nb, Fe a n d / o r enrichment of C more than 0.03
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Fig. 9. Relation between experimental and predicted oxygen pickup (OP) after exposure to HTGR helium environment at 1000°C for 1000 h. Predicted OP=0.8269+0.5684Mn1.9013Si-74.8233B-53.8198¥+3.3176Ti where OP is in mg/cm 2, and Mn, Si, B, Y and Ti are in mass%.
311
Y. Muto etal. / High temperature structural materials for HTGR Creep rupture strength
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Fig. 10. Comparison of newly developed superalloy (predicted values) with the existing high-temperature alloys. mass% exhibits no beneficial influence and that the addition of B and Y in either singular or simultaneous forms is remarkably efficient in improving either creep rupture strength, tensile ductility, corrosion resistance or hot workability. As for the effect of Mn, Si and Ti, they have both beneficial and detrimental effects. Figure 9 is an example of multiple regression analysis results that shows the relation between experimental and predicted oxygen pickup after exposure to a helium environment at 1000°C for 1000 h [19]. Based on the experimental results, it has been judged that Ni/18 to 19 mass% Cr/20 to 22 mass% W/0.03 mass% C/0.08 mass% Ti/0.02 to 0.05 mass% Zr/0.002 to 0.007 mass% Y/0.0035 to 0.006 mass% B is the optimum chemical composition. And the newly proposed N i - C r - W superalloy can be expected to in-
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dicate well-balanced properties and superior performance to the existing high-temperature alloys as illustrated in Fig. 10.
4. Structural test for the high temperature design code In order to develop an advanced life prediction method at the high temperature exceeding around 800°C, developing both a constitutive equation and an improved creep-fatigue damage evaluation method is required. Then, efforts are focussed on it. 4.1. Constitutive equation
Constant stress-creep tests were conducted at 900°C for Hastelloy XR2 which is a modified version of HastelloyXR. A uni-axial creep constitutive equation of the Garofalo type was made which had three parameters such as Et: asymptotic primary strain, r: recipro-
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Fig. 13. Prediction by the time fraction rule.
312
Y. Muto et al. / High temperature structural materials for HTGR r
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Fig. 14. Prediction by the ductility exhaustion rule. cal of time constant and groin: minimum creep strain rate [20]. The values of the parameters were determined experimentally. The parameters r and Emin were fitted very well, though a significant data-scattering was observed in the parameter Et. Therefore, it can be said that this type of creep equation must be one of the most adequate ones for Hastelloy XR2. The equation was applied to predict a creep behavior under a variable stress condition using the strain hardening law. The result was, however, insufficient and suggested the effect of an internal back stress. Then, a new constitutive equation incorporating the internal back stress was made where a creep strain rate equation was assumed to be the well-known power law type [21], and an evolution equation of the internal back stress was based on a creep-recovery model [22]. The parameters were determined by stress dip tests whose results are shown in Fig. 11. The new constitutive equation is given as follows [23]. - ~ b ( ~ ) ]~('~',
(5)
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(6)
= A(~)[~
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300
400
500
600
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where A 0 r ) , B(~r) and D((r) are material constants, (rb. 0 = an initial value of the internal back stress. The creep behaviors were calculated by both the
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Fig. 17. Cyclic thermal loading of cylindrical specimen.
Y. Muto et al. / High temperature structural materials for HTGR
new equation and the Garofalo equation for variable stress loads. It was ascertained that the new constitutive equation could predict the creep behavior under the variable stress loads more exactly than the Garofalo equation as shown in Fig. 12. A stress-relaxation behavior is, however, frequently observed in thermal loading cycles and a creep damage during the relaxation becomes important. Therefore, further work is needed to examine the validity of the new creep constitutive equation for the relaxation. 4.2. Creep-fatigue tests
The low cycle fatigue tests with and without a dwell period were conducted at 850°C and 950°C under vacuum conditions of 10 3 Pa. The strain rate in loading and unloading is changed from 0.01%/s to 0 . 5 % / s at 850°C, but to l % / s at 950°C. The dwell period in a cycle ranges from 150 s to 1800 s at tension a n d / o r compression sides: this time duration is enough to achieve the low level of residual stress rate where the cycles to failure won't be changed. Cycles to failures were predicted by the time fraction rule [24], a ductility exhaustion rule [25] and a damage rate equation [26], whose results are shown in Figs. 13, 14 and 15, respectively. In the time fraction rule the predicted cycles to failure were larger than the real ones in TH and CH
313
cycles: TH is a cycle with a dwell period in tension and CH in compression. Consequently a failure will occur at the creep-fatigue damage less than unity in these cycles. The ductility exhaustion rule could take the effect of strain rate into account. The method is more likely to reflect the mechanism of creep damage rather than the time fraction rule though the accuracy of preditions by both methods are the same. It is, however, still suggested from the results that the damage is too much counted in TCH cycle: TCH is a symmetrical cycle with a dwell period in tension and compression sides. A concept of the damage-rate equation is basically the growth of cavities, described as a function of inelastic strain range. The damage or the cavity growth in the CH cycle will need further discussion. 4.3. Biaxial test
Biaxial stress-rupture tests were conducted for the condition of both the combined tension and torsion loads and the combined compression and torsion loads. It was ascertained that a uniaxial tensile stress-rupture test results in the shortest rupture life for the same equivalent stress of constant loading conditons [12]. Then, tests were succeeded under constant equivalent stress conditions. Thin-walled cylindrical test specimen of 12 mm gauge length, 22 mm outer diameter and 2 mm wall thickness were used.
Fig. 18. Cyclic bending of curved tube.
314
Y. Muto et aL / High temperature structural materials fi)r HT(;R
Creep strain behaviors are shown in Fig. 16 by a parameter of angle in the stress space. The creep strain rates in the primary creep region is retarded as the angle decreased. The data of creep strain exceeding several percents are not reliable because those were accelerated by buckling. Buckling due to the eccentricity of the loading axis and bulging type bucklings were observed in the case of torsion and compression, respectively. Biaxial creep-fatigue tests with constant torsion and cyclic axial load were conducted at 900°C. The experimental creep-fatigue life coincided with the analyrical one predicted by the time fraction rule. 4.4. Component test
a thick wall will take thermally cycled loading and a heat transfer tube will take mechanically cycled loading. Photos of the former test section and the latter test apparatus are shown in Figs. 17 and 18, respectively. In curved tube-cyclic bending tests at 900°C, cracks were observed to initiate at the bottom of the curved corner and run to the circumferential direction in the case of in-plane bending tests, while they initiated at the side and propagated to the direction of 45 ° to the tube axis in the case of out-of-plane bending test. The results coincide with those by the stress analyses as shown in Fig. 19. 5. Conclusions
The limits of thermal loading to the progressive time-dependent damage of high temperature components in the intermediate heat exchanger system have heen ~tndied in mnlti-n~zial ,~tre,:~ ~t~te~" a cylinder with
(1) For 2 l / 4 C r - l M o steel the fracture toughness at 400°C, the service temperature of the pressure vessel, was estimated to be = 70 k J / m z after irradiation, ten times as high as H T T R life fluence.
/
I Maximumshear stress {n-plane bending \ (Out-of-planebending~/~ (/ (Experiment) Out-of-planebending ><~ ,/VI '~ ..... (Experiment) ~,.~. ~ ,) Loaa,ng s,ae / [
" [ \-.. I Maximumlongitudinalstress / ~J (In-planebending(Analysis)) Fixed side Fig. 19. Appearance of failure tubes.
Y. Muto et al. / High temperature structural materials )br HTGR (2) T h e d e c r e a s e in the fracture t o u g h n e s s of thermally aged material caused by irradiation is m u c h less t h a n that of the as-received base plate material. (3) T h e radiation h a r d e n i n g b e c a m e very small at temp e r a t u r e s higher than 350°C with a n e u t r o n fluence of a b o u t 3 × 1023 n / m 2 ( E > 1 MeV). (4) Hastelloy X R base metals as well as filler metals for the weldings have b e e n developed and characterized to apply to the high t e m p e r a t u r e compon e n t s of the H T T R . (5) O p t i m u m chemical composition of newly developed N i - C r - W superalloy, which can be expected to indicate well-balanced properties, has b e e n proposed. (6) T h e constitutive e q u a t i o n incorporating the internal back stress could r e p r e s e n t the creep strain b e h a v i o r of Hastelloy X R at variable stress conditions. (7) Applicabilities of the time fraction rule, the ductility exhaustion rule and the d a m a g e rate e q u a t i o n to Hastelloy X R were e x a m i n e d b a s e d on experim e n t a l data. (8) Biaxial c r e e p a n d creep-fatigue data were o b t a i n e d for Hastelloy XR.
References [1] [2] [3] [4]
S. Saito et al., Nucl. Engrg. Des. 132 (1991) 85. T. Kodaira et al., JAERI-M 85-170 (1985). T. Oku et al., Nucl. Engrg. Des. 119 (1990) 177. Y. Nishiyama et al., ASTM STP 1125, R.E. Stroller et al., Eds. (1992) pp. 1287-1303. [5] W.A. Logsdon and J.A. Begley, Eng. Fracture Mechs. 9 (1977) 461.
315
[6] T. Iwadate et al., ASTM STP 631, J.M. Barsom, Eds. (1977) pp. 493-506. [7] M. Suzuki et al., ASTM STP 956, F.A. Garner et al., Eds. (1987) pp. 98-110. [8] M. Shindo and T. Kondo, Proc. Int. Conf. on Gas-Cooled Reactor Today, Bristol, UK, 1982 (British Nuclear Energy Society) 2, pp.179-184. [9] Y. Kurata et al., Proc. Int. Conf. on Creep, Tokyo, Japan, 1986 (The Japan Society of Mechanical Engineers), pp. 97-102. [10] K. Hada et al., JAERI-M 911-148 (1990). [11] T. Tsukada et al., Proc. Workshop of Advanced Materials and Protective Coatings, Tokyo, Japan, 1992 (Elsevier Science Publishers) pp. 233-242. [12] K.Hada et al., Nucl. Engrg. Des. 132 (1991) 1-11. [13] H. Tsuji and T. Kondo, J. Nucl. Mater. 151t (1987) 259265. [14] H. Tsuji and H. Nakajima, J. Nucl. Mater. 158 (1988) 267-278. [15] H. Tsuji and H. Nakajima, J. Nucl. Mater. 151 (1987) 1-9. [16] T. Udoguchi and T. Nakanishi, Int. J. Press. Vessel and Piping 9 (19811 107-123. [17] K. Watanabe et al., J. Nucl. Mater. 185 (19911 8-18. [18] Ad Hoc Committee on Advanced Superalloys, Technical Expert Committee on HTGR, Japan Atomic Energy Research Institute, JAERI-M 88-270 (1989). [19] H. Tsuji et al., JAERI-M 91-136 (19911. [20] F.Garofalo, Fundamentals of Creep and Creep-Rupture in Metals (Macmillan, New York, 1965). [21] F.H. Norton, Creep of Steel at High Temperatures (McGraw Hill, New York, 1929). [22] R. Lagneborg, Metal Science J. 3 (1969) 161-168. [23] Y. Kaji et al., The 5th Int. Conf. on Creep of Materials (1992). [24] Cases of ASME Boiler and Pressure Vessel Code, Case N-47. [25] R.H. Priest and E.G. Ellison, Material Science and Engrg. 49 (19811 7-17. [26] S. Majumdar and P.S. Maiya, J. Engrg. Materials & Technol. 102 (1980) 159-167.