Mechanical property characterization in support of elevated temperature reactor design

Nuclear Engineering and Design 45 (1978) 225-242 © North-Holland Publishing Company

225

MECHANICAL PROPERTY CHARACTERIZATION IN SUPPORT OF ELEVATED TEMPERATURE REACTOR DESIGN * R.L. KLUEH, C.R. BRINK_MAN and V.K. SIKKA Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37830, USA

Received 29 April 1977

The mechanical properties data used for the design of nuclear reactors and reactor components for operation at elevated temperatures often show considerable variation. In types 304 and 316 stainless steels, much of the heat-to-heat variation is due to variation in chemical composition and grain size. For a heat-treatable ferritic steel such as 2 1/4 Cr-1 Mo steel, chemical composition is again important, but the heat treatment is more important. An understanding of the metallurgical processes that lead to data variation and property changes in these steels can be used to optimize reactor design.

1. Introduction To develop the constitutive equations that designers require for materials to be used for the design of nuclear reactors and reactor components for elevatedtemperature service, substantial amounts of mechanical properties data and behavioral studies are required. In addition to this requirement, designers also need absolute reference values for mechanical properties over the stress and temperature range of interest for the design. Because of such requirements, not only must accurate mechanical properties data be obtained by laboratory measurements, but also they must be collated, correlated, and presented in a form that can be used by the design community. Such design data are now being generated for materials to be used in liquid-metal fast breeder reac. tors: type 304 stainless steel (piping and structural material), type 316 stainless steel (primary vessel and intermediate heat exchanger material), and 2 1/4 Cr1 Mo steel (steam generator material). Stainless steel properties are being determined over the range 25 to 649°C and 2 1/4 Cr-1 Mo steel properties over the range 25 to 593°C. Data have been or are being ob* Research sponsored by the Engery Research and Development Administration under contract with the Union Carbi~te Corporation.

tained to define the following uniaxial properties: tensile, creep [ 5 - 7 ] , relaxation [8], fatigue [9-11], creep-fatigue interactions [ 10,11 ], and subcritical crack growth [10,11]. Also, load-history tests are being made. The load-history tests are required as a basis for verifying inelastic constitutive equations and are used in conjunction with the constant load creep data to determine the hardening laws. These tests inelude uniaxial creep with periodic load changes and creep tests under stepwise-alternating (fully reversed) stresses. Some uniaxial creep data generated for annealed 2 1/4 Cr-1 Mo steel at 566°C are shown in fig. 1. Some results of an experimental creep test with load changes are shown in fig. 2, along with a strain-hardening prediction [12] that was determined using the curves in fig. 1. Two creep tests (fig. 3) with stress changes gave data that supported a strain-hardening law for annealed 2 1/4 Cr-1 Mo steel [12]. The examples of mechanical properties data (figs. 1, 2 and 3) were obtained for one heat of steel for a given heat treatment. However, the mechanical properties for a given structural material can show appreciable variation from many causes: heat-to-heat variations in chemical composition or grain size and variations in heat treatment and processing; experimental variations among different investigators can also be important. RegardleSs of the source of variation, the

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designer requires a knowledge of that variation so that proper safety factors may be included. The ASME code body sets allowable design stresses based on minimum values obtained from a statistical treatment of the data. Likewise, in the Nuclear Systems Materials Handbook [13], for a given mechanical property, expected and minimum values are statistically determined from the range of experimental values obtained from various heats by numerous experimenters. Although such a statistical treatment of the data is adequate to meet the needs of the designer, an understanding of the underlying causes for that variation would be extremely useful. If the reasons for the mechanical properties variation were understood, it should thenbe possible to produce alloys with the desired strength properties and properties that fall within a much narrower scatterband than presently observed. Furthermore, with an understanding of the effects on mechanical properties of such things as

chemical composition, grain size, heat treatment, etc., the designer could more accurately represent the true characteristics of the actual heat(s) of material used for a given reactor or reactor component. At present the Mechanical Properties Group in the Metals and Ceramics Division at Oak Ridge National Laboratory is engaged in the following studies: (1) development of experimental design data, (2) data analysis for inclusion in the ASME codes and Nuclear Systems Materials Handbook [13] to be used for reactor design, and (3) determination of the effects of various processing variables on the mechanical properties. The objective of this paper is to review some of the types of design data that have been generated, show the variations that result when all the available data (from various sources) are collected, and show how these data are statistically analyzed. Finally, we intend to review some of the results of

228

R.L. Klueh et al. / Elevated temperature reactor design

our studies that have been undertaken to understand the causes of mechanical property variation and to show how an understanding of the metallurgical causes behind these variations can be of use to the designer. Since we cannot possibly discuss all mechanical properties and their variation, most of the discussion will be concerned with uniaxial tensile and creep properties, with a brief discussion on fatigue properties. However, the principles outlined for the understanding of the tensile and creep properties should also apply for the other mechanical properties.

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For elevated temperature design, tensile and creep properties data are used to set the design allowable stress limits. For ASME code case 1592, the 0.2% yield and ultimate tensile strength minimum values set the time independent stress limit Sin, while the stress to cause rupture, the stress to cause tertiary creep in a given time, and the stress to give 1% strain in l0 s h determine the time dependent stress intensity limit S t . The evaluation of these allowable stresses depends on the availability of data, variations in the data, and the methods for analysis. For the present paper, the data variations are of most interest. In this section o f the paper, some examples o f the data variation for the ferritic 2 1/4 Cr-1 Mo steel and the austenitic stainless steels types 304 and 316 will be given.

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Annealed 2 1/4 Cr-1 Mo steel, which is to be used as the structural material for the steam generator of the Clinch River Breeder Reactor, shows quite large variations in properties [ 14,15]. Examples of the behavior encountered are shown in figs. 4 and 5. To approximate these data for design use, they are statistically analyzed and expected values are determined, along with upper and lower tolerance limits [16] (figs. 4 and 5). The tolerance limits are calculated such that at a confidence level of 0.95, 90% of all observed values are expected to fall within the limits (95% are expected to be above the lower tolerance limit). For handbook use, the ultimate tensile strength was expressed [16] as a cubic polynomial

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229

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in temperature. Su = 527 - 1.03T+ 4.54 × 10-aT 2

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where Su is the tensile strength in MPa and T is the temperature in °C. A similar procedure was used for other tensile properties [15]. By a parametric technique, the expected rupture life was found to be related to stress and temperature from 371 to 593°C according to, log tR = -31.45 + 3.93 × 1 0 4 / T - 19.75 log o ,

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where tR is the time to rupture in h, T is the temperature in K, and o is the stress in MPa. Although relationships of the type given above can be used as an "average expected value" for these properties, these relationships are strictly empirical and lead to no understanding of the variation in properties. Furthermore, data for creep rupture are only available out to about 20,000 h at 538°C (fig. 4). At some of the lower temperatures, data were not available beyond about 5000 h. Nevertheless, reactor applications require extrapolations for lifetimes of 3 × l0 s h. Of course, extrapolations can be made with eq. (2). However, only with some understanding of the metallurgical effects of long-time exposure at elevated temperatures, can such extrapolations be made with any confidence. These metallurgical effects will be discussed in sect. 3 of this paper. Z Z Austenitic stainless steels types 304 and 316 Gross variations in tensile and creep properties similar to those observed for 2 1/4 Cr-1 Me steel are also observed for the austenitic stainless steels [17-20]. Techniques similar to those already described (sect. 2.1) can be used to represent the property variations. Recently, available data from around the world were examined, and similar variations were found, regardless of the country of origin [21]. The variations are shown when the yield stress and ultimate tensile strength are presented as functions of temperature for data obtained in Britain, Japan, and the United States (fig. 6). When the strengths and variations for different product forms were compared, it was found that bar

product made in the United States is stronger than the tube and plate products, whereas the bar product made in Britain and Japan are weaker than the other products. The reasons for such differences in strength are probably associated with the difference in fabrication sequences or the final finishing treatment in the different countries. That is, different amounts of cold work are present in the different products. This difference in strengths due to varying amounts of cold work in the different product forms probably has a large effect on the variability (fig. 6). To evaluate the effect of cold work, the differences in tensile strength for mill-annealed (as received) and laboratoryannealed (reannealed) heats were examined for various product forms of types 304 and 316 stainless steels (a 0.5 h anneal at 1065°C effectively removes the cold work). These studies confirmed the results stated above concerning the product forms. It was found that the bar product for the austenitic stainless steels (from U.S. vendors) contained about 10% cold work compared to 3-4% for the plate and pipe [221. The effect of cold work was also examined for the creep-rupture properties [5]. Twenty different heats of type 304 stainless steel and seven heats of type 316 stainless steel were examined in the mill-annealed (as-received) and reannealed condition. All heats met appropriate product specifications as set by ASTM. All heats in both the as-received and reannealed conditions met the ASME code minimum values. Large variations were observed in creep-rupture properties of the heats of type 304 stainless steel at all test temperatures. Reannealing the as-received material, which should remove cold work, lowered the rupture time, but variations were still observed. For example, at 593°C and 207 MPa, the following ranges of creep properties were observed for the twenty heats of type 304 stainless steel in the reannealed condition: Rupture Life Minimum Creep Rate Loading Strain Creep Strain Reduction in Area

84--2580 h, 0.00077-0.16%/h, 4.8-8.8%, 3.6-45.5%, 11.5-57.7%.

Similar variations were also found for type 316 stainless steel [5].

230

R.L. Klueh et aL / Elevated temperature reactor design

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3. Metallurgical basis for variation As explained in the previous section, the mechanical properties of the austenitic stainless steels type 304 and type 316 and the ferritic 2 1/4 Cr-1 Mo steel show large variations in mechanical properties. The reasons for these variations, usually called heat-to-heat variations, are many. Designers often assume that, regardless of the alloy, the underlying causes for such variations are similar; the following discussion shows that this is not the case. Although differences in properties result from differences in processing, the underlying causes for these differences are not the same for the austenitic and ferritic materials. Whereas the heat-to-heat differences

in the austenitic stainless steels are primarily due to variations in chemical composition and grain size, differences for the ferritic steels result from the heat treatment they undergo. During the heat treatment, which is the basis for the strength of these steels, 2 1/4 Cr-1 Mo steel is austenitized (transformed to face-centered cubic austenite) by heating above the Ac3 temperature. Then on cooling, the austenite transforms back to body-centered cubic ferrite; several constituents occur in the ferritic condition, depending on the heat treatment procedure. It is this transformation that causes the major property variations noted in ferritic steels. In the following sections, the metallurgical reasons for the property variations in the ferritic 2 1/4 Cr1 Mo steel and the austenitic stainless steels types 304 and 316 will be discussed separately. 3.1. Ferritic 2 1/4 Cr-1 Mo steel

The basis for the strength of a heat treatable steel such as 2 1/4 Cr-1 Mo steel is the austenite-ferrite

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transformation. This transformation can be most easily understood with reference to an isothermaltransformation diagram or a continuous-cooling transformation (CCT) diagram. A CCT diagram for 2 1[4 Cr1 Mo steel is shown in fig. 7 [23]. The first step in heat treating any ferritic steel is to austenitize it (transform the ferrite to austenite) by heating it above the Ac3 temperature (it is usually held at this temperature about 1 h per inch of section thickness); the A 3 is about 900°C for 2 1/4 Cr-1 Mo steel (fig. 7). The steel is then cooled by one of several standard techniques (heat treatments): quench (in water or oil), normalize (air cool), or anneal. For a given heat treatment, the microstructure that evolves depends on how rapidly the steel is cooled. To determine the constituents that evolve in the microstructure, the cooling curve for a given heat treatment can be superimposed on the diagram in fig. 7. The amounts of the constituents (i.e., proeutectoid ferrite, bainite,

etc.,) that evolve will depend on the time the steel is at the temperature of that respective region as shown in the CCT diagram. Quenching and normalizing are rapid cools. Depending on the section size that is heat treated, large amounts of bainite can form. Bainite has very high strength with limited ductility, which can be improved by tempering. Therefore, these steels are generally used after tempering, although tempering also lowers the strength somewhat. The section size effect illustrates one of the inadequacies of labeling the heat treatment by the cooling procedure. The properties of the steel are determined by the microstructural constituents. For a large normalized and tempered forging, for instance, the microstructure may contain both bainite and proeutectoid ferrite, with the relative amounts of these constituents varying through the cross section (i.e., the outside cools faster than the inside and will contain more bainite). On the

232

R.L. Klueh et al. / Elevated temperature reactor design

other hand, a normalized and tempered 1-in. plate or rod may be completely bainite [6]. The terms, "annealed" or full anneal are used rather loosely by the steel industry [24]. A steel is annealed by slowly cooling (in a furnace) from the austenitizing temperature to cause the transformation of the austenite to a relatively soft ferrite-pearlite aggregate. That is, most of the transformation occurs when the steel is being cooled through the proeutectoid ferrite region (fig. 7). Generally, a steel is annealed by furnace cooling from the austenitzing temperature. For tubing, however, an isothermal anneal is generally used to obtain the same ferrite-pearlite aggregate. An isothermal anneal is one in which, after austenitizing, the steel is cooled to some intermediate temperature in the proeutectoid ferrite transformation region as shown in the CCT diagram (fig. 7), held at this temperature 2 - 4 h, and then cooled to room temperature. For tubing this can be accomplished in a continuous annealing furnace that contains different temperature zones. The tubes are continuously moved through the furnace, from a zone that is maintained at the austenitizing temperature into a zone that is held at a constant temperature, usually between 677 and 732°C. Steels given either a full anneal (continuously furnace cooled) or an isothermal anneal are generally referred to simply as annealed. Since the steam generator of the Clinch River Breeder Reactor is to be constructed of annealed 2 1/4 Cr-1 Mo steel, most of our present studies are on that heat treated condition, and the following discussion will be directed to the steel in that condition. Actually, the CCT diagram (fig. 7) is directly applicable only for the steel for which it was determined, which is why the data on chemical composition and grain size are listed at the top of the figure. Heat-toheat variations in chemical composition can lead to changes in the CCT diagram. The primary effect of small compositional changes is to move the start of the proeutectoid ferrite transformation region to longer or shorter times, thus leading to more or less bainite in the microstructure for a given cooling rate [25]. However, most annealed 2 1/4 Cr-1 Mo steel is cooled slowly enough to form a microstructure that is primarily proeutectoid ferrite, with only small amounts of pearlite and bainite (usually less than 15-20%). As discussed below, the difference in properties is therefore primarily the result of differences

in the proeutectoid ferrite [26]. One deficiency of CCT diagrams (or isothermal transformation diagrams) should be pointed out: Although proeutectoid ferrite and pearlite regions are shown (fig. 7), no account is taken of the fact that MozC precipitate eventually forms in the proeutectoid ferrite [27]. This precipitate can form during the heat treatment (depending on the cooling rate) or while the material is in test or service. Later in this paper, the important role that the Mo2C formation rate plays in the effect of heat treatment on the mechanical properties of the annealed steel will be shown. Studies have been completed to determine the effect of heat treatment on the tensile [26] and creeprupture [7] properties of 2 1/4 Cr-1 Mo steel to show how heat treatment effects can lead to variations like those shown in figs. 4 and 5. Sections of a 25 mm (1 in.) plate from a single heat were put in the annealed condition by austenitizing for 1 h at 927°C and then furnace cooling. A section was also isothermally annealed by furnace cooling to 704°C at about 83°C/h, held at that temperature for 2 h, then furnace cooled. The cooling curves for these three heat treat conditions are shown in fig. 8. Hereafter, the two annealed plates will be referred to as AN-1 and AN-2 (AN-2 was cooled faster than AN-I) and the isothermally annealed plate as IA. According to the CCT diagram (fig. 7), most of the transformation to proeutectoid ferrite occurs at temperatures above about 650°C. Both AN-1 and IA reach this temperature in about six hours (an average cooling rate to 650°C of about 46°C/h), while the AN-2 reaches this temperature in about 5 h (an average cooling rate of about 55°C/h). The microstructures of the steel in these three heat treated conditions were primarily proeutectoid ferrite with some peralite and bainite. Heat treatments AN-l, AN-2, and IA were estimated to have respectively 15, 20, and 15% bainite and small amounts of pearlite [26]. Both tensile and creep-rupture properties were determined for these three different heat treated conditions. Tensile tests were made between 25 and 593°C at strain rates of 2.67 × 10 -6 , 6.67 × 10 -s, 6.67 × 10 -4, and 6.67 × 10-3/s. Over this range of test variables, there were significant differences in properties, which had to be attributed to the heat treatment. The yield

R.L. Klueh et aL / Elevated temperature reactor design

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strengths of AN- 1 and IA showed very little difference, while AN-2 was slightly stronger (fig. 9). AN-2 also showed much less of a decrease in yield strength with increasing temperature and showed essentially no effect of strain rate, whereas, the other *-vo heat

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234

R.L. Klueh et aL / Elevated temperature reactor design

1 Mo steel gathered from published literature and industrial laboratories [14]. For the tensile properties of the annealed steel, Smith gave a scatterband of data: all three of our heat treatments give yield stresses that fall near the center of Smith's scatterband (fig. 11). Only at the highest test temperatures do the properties of AN-2 deviate from the center of the scatterband into the upper half of the scatterband; this deviation occurs for AN-2 because the yield stress for this material does not decrease as rapidly with temperature as the yield stress does for the other two heat treatments. The ultimate tensile strength behavior was somewhat more complicated than the yield stress behavior. In fig. 10, the ultimate tensile strengths for the three heat treatments are shown as a function of temperature for strain rates of 2.67 × 10 -6 and 6.67 × 10-3/s. Again, AN-1 and IA show similarities: the ultimate tensile strengths are the same between 25 and 200°C at both strain rates; they are also the same from 510 to 593°C at 2.67 × 10-6/s and from about 371 to 593°C at 6.67 × 10-3/s. At the intermediate temperatures, (where the peaks occur in the curves of fig. 10),

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600

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there are significant differences betweenthe IA and AN-1 heat treatments and AN-2. In the regions of the peaks (fig. 10), dynamic strain aging occurs, and it is in the range of these strain aging peaks that the largest deviations for the three heat treatments occur. At 2.67 × 10-6/s, the strain aging peaks for AN-1 and IA occur at about the same temperature - near 315°C - while the peak for AN-2 is shifted to a somewhat higher temperature - about 371°C. This shift to a higher temperature for AN-2 is also obvious at 6.67 × 10-3/s. When the data (fig. 10) are compared with the scatterband from Smith's compilation for the annealed steel, the curves for all three heat treatments at a strain rate of 6.67 × 10-3/s are found to fall within the scatterband. Although Smith gives no details on the strain rates of the tests in his compilation [14], it is assumed that most of the tests conformed to the ASTM specification E-151; the ultimate tensile strength in most of those tests would probably be reached at a strain rate of about 6.67 × 10-3/s or greater [E- 151 recommends a strain rate of 0.005/min (8.33 × 10-S/s) to the 0.6% offset strain and 0.l/min (1.67 × 10-3/s) to failure]. Thus, the ultimate tensile strength data for all three heat treatments tested at typical strain rates would appear to fall within the scatterband given by Smith. Heat treatment can likewise affect the creep-rupture properties of 2 1/4 Cr-1 Mo steel and produce strength variations (fig. 5) [7]. At high stresses there is an obvious difference between the properties of AN-2 - the fast-cool anneal - and the other two heat treatments (fig. 11). The properties for AN-1 are

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R.L. Klueh et al. / Elevated temperature reactor design

slightly above those for IA, but as the stress is decreased, the rupture times for these two steels also approach one another. At 566°C the differences in the creep-rupture properties of the three heat treatments were not nearly as great and approached each other in much shorter times. At 454°C, where only AN-1 and IA were tested, the properties for AN-1 were significantly greater than those for IA, and this difference persisted beyond the test times used in the study [7]. Explanations for the effect of heat treatment are all intimately connected with the microstructure developed during the heat treatment and the continuous change that the microstructure undergoes during elevated-temperature exposure. It was established that for these annealing treatments the properties were determined by the proeutectoid ferrite microstructure, and the greater amount of bainite present in the AN-2 heat treatment had only a minor effect [7,26]. Although the proeutectoid ferrite contains Mo2C precipitate, the above observations on tensile and creep-rupture properties (also on fatigue properties, to be briefly discussed in section 4) are explained as effects of alloying elements in solution. On an atomic scale, plastic deformation is the result of dislocation motion though the atomic lattice of the alloy. Dislocations, which are line defects that separate slipped from unslipped material, move in response to an imposed stress. Any process that hinders dislocation motion through the alloy lattice strengthens the alloy. Since the metal lattice is distorted in the vicinity of a dislocation, it is a highenergy region in the lattice. The energy can be lowered by attracting an atmosphere of alloying elements (i.e., large solute atoms will segregate to the expanded region of the lattice, the small interstitial atoms to the compressed region). Such an atmosphere can then hinder dislocation motion. That is, for plastic deformation to proceed by the movement of atmospherecontaining dislocations, the dislocations must either be pulled from the atmosphere or the atmosphere must move with the dislocation by atomic diffusion. In 2 1/4 Cr-1 Mo steel, the dynamic strain-aging peaks (fig. 10) for the ultimate tensile strength-temperature relationships are the result of interaction solid solution hardening [28,29]. Interaction solid solution hardening is the result of the formation of dislocation atmospheres by carbon and molybdenum

235

atom clusters that impede dislocation motion. This interaction of Mo-C clusters with dislocations leads to alloy strengthening, which in turn leads to those properties discussed above. The dynamic strain-aging peaks (fig. 10) result because it is at the temperatures of the peaks that the Mo-C atmosphere clusters have enough mobility to rather quickly form a dislocation atmosphere when the dislocation sweeps the lattice. The dislocations are thus immobilized, leading to a strength peak. However, as the temperature is increased the atoms that form the dislocation atmosphere acquire enough mobility to move with the dislocation by diffusion, thus losing their strengthening effect (i.e., dislocation motion is not impeded and the strength drops off). These ideas explain the effect of the heat treatment on the dynamic strain-aging peak. During the anneal, Mo2C precipitates in the proeuteetoid ferrite matrix. For precipitation to proceed, diffusion is required. Since diffusion requires relatively high temperatures, the slower the furnace cool during the anneal heat treatment, the more Mo2C that is formed. And the more Mo2C formed, the less molybdenum and carbon that remain in solution and are thus available to form dislocation atmospheres. Since the AN-2 heat treatment had the fastest cool, it had the most molybdenum and carbon remaining in solution that could then enhance the strength by interaction solid solution hardening. Thus, the highest strength peak is observed for this steel. A similar explanation applies for the enhanced creep-rupture strength of AN-2 at high stresses (short rupture times) (fig. 13) [29]. Since more molybdenum and carbon remain in solution for this heat treatment, initially AN-2 has the highest strength. However, at the elevated temperature of the creep test, precipitation of Mo2C continues. Eventually, the molybdenum and carbon concentrations in solution of all three heat treatments fall to values where further strengthening due to interaction solid-solution hardening is negligible. When this happens, strengths for all three heat treatments become similar so that with decreashag stress, the strengths tend to become identical. The above discussion illustrates the importance of metallurgical studies in conjunction with the determination of design properties for 2 1/4 Cr-1 Mo steel. Based on the above discussion, five other examples of heat treatment effects come readily to mind:

R.L. Klueh et aL / Elevated temperature reactor design

236

1. Since the dynamic strain-aging peak is caused by molybdenum and carbon atoms in solution, the peak will eventually disappear after elevated-temperature exposure as the molybdenum and carbon are removed from solution by precipitation. 2. Since the high cycle fatigue properties are in large measure determined by the tensile properties, the deterioration of properties in the dynamic strainaging region should be considered by the designer. 3. That the creep-rupture properties for the three different heat treatments approach each other with time indicates that the same approach of properties should also occur when different heats are tested, since with time at temperature, for all cases the microstructure becomes a ferrite matrix containing large carbide precipitates [27]. This fact may, in turn, aid in making long-time extrapolations. 4. The effect of thermal aging on the precipitate particles further affects properties. The Mo2C carbides

that form are unstable and give way to M6C, which reduces the strength [30]. 5. In many instances the creep curves for 2 I/4 Cr1 Mo steel are nonclassical. Instead of having single primary, secondary (steady state), and tertiary creep stages, two steady state stages are observed. The metallurgical explanation for this behavior also involves interaction solid solution hardening, and the explanation has been used to determine proper design values [7]. 3.2. Austenitic stainless steels type 304 and type 316

Since the austenitic stainless steels do not undergo a phase transformation once solidified, the metallurgical basis for the heat-to-heat variations must be inherent to the austenitic structure. For this discussion the metallurgical effects [e.g., chemical composition, precipitation, and grain size) will be separated from

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R.L. Klueh et al. / Elevated temperature reactor design

237

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cold work effects which were discussed in the previous section (sect. 2.b)]. Hence, only data from reannealed specimens (0.5 h at 1065°C, which removes the cold work) will be discussed. As pointed out in the previous section, large heatto-heat variations still remain after the cold work is removed by reannealing. A review of the literature shows that interstitial atoms such as carbon and nitrogen are important strengthening elements in steel. Furthermore, microstructural variations such as grain size are particularly important in determining strength properties. For both types 304 and 316 stainless steel, the strength properties were found to be directly proportional to the total interstitial concentration and inversely proportional to the square root of the grain size [51:

where S is the yield stress or ultimate tensile strength, C and N are the concentrations of carbon and nitrogen respectively, D is the average grain diameter determined by the line intercept method, and So and k are constants. In fig. 12, the ultimate tensile strength data at four different temperatures are plotted as functions of (C + N ) D - 1 / 2 for nineteen heats of type 304 stainless steel, and fig. 13 shows the plot for seven heats of type 316 stainless steel. The least squares lines show excellent fit to the data. It has also been shown that the creep-rupture properties of types 304 and 316 stainless steel can be related to the tensile properties by the relationship [5],

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238

R.L. Klueh et al. / E l e v a t e d temperature reactor design

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where o is the stress to cause rupture in time tR and a2 and et3 are constants [/3 and Su are the same as for eq. (4)]. These relationships can account for the variations (fig. 14). Fits of the data with and without the use of the tensile strength correlation are given; the standard error of estimate in log tR decreased from 0.52 to 0.39 at 593°C and from 0.49 to 0.34 at 649°C when the ultimate tensile strength of the heat was taken into account. By combining eqs. (3) and (5), a relationship between rupture life and interstitial content and grain size result. However, when creep data were plotted as a function of (C +N) D -1/2, it was found that there is a distinct branching in the curves for weak and strong heats. Although (C +N)D -112 accounts for both sets, an additional factor is required to bring strong heats in line with the weak heats. A detailed

chemical analysis indicates that the residual niobium content is at least partly responsible for these strength differences. The general trend indicated that heats containing 26, 100, and 160 ppm correspond to weak, intermediate, and strong heats, respectively (these are average niobium contents for several heats). Work is now in progress to check the influence of niobium on special heats prepared at known niobium levels.

4. Other mechanical properties The previous two sections were directed toward tensile and creep properties. Metallurgical effects on other properties must also be considered. Fatigue is one area where a limited amount of work has been done.

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Since LMFBR components will be subject to many temperature changes, and hence cyclic stresses, design against possible fatigue-induced failure becomes an important consideration. ASME Code Case 1592 employs the results of uniaxial fully reversed strain and load-controlled fatigue tests to establish the allowable number of cycles that a component can be subjected to without danger of fatigue-induced failure. These design allowable cycles are established using large safety factors and are thought to be so conservative as to preclude crack growth. A typical set of uniaxial isothermal fatigue curves with experimental data for 2 1/4 Cr-1 Mo steel at several temperatures of interest is shown in fig. 15. The data points were generated at several different laboratories on approximately six different heats of

material which conformed to the ASME specifications SA-336-F22A and SA-387-22C1. Comparing the best-fit lines in fig. 15 for the three different temperature ranges shown, it is apparent that increasing the temperature decreases the cyclic life for a given strain range. Further, considering the scatter inherent in fatigue test data and considering that the test results were generated from several heats of material, no heat-to-heat variations or heat treatment effects (i.e., isothermal versus annealed) are apparent. Ellis et al. [31 ] compared results from a single heat of 2 1/4 Cr-1 Mo steel that had been given either a full anneal or an isothermal anneal. They found that in terms of cyclic life the test results were nearly identical. However, the cyclic stress response for a given

240

R.L. Klueh et al. / Elevated temperature reactor design

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Fig. 16. Cyclic stress range at half the cycle life (Nf/2) as a function of temperature at total strain ranges (Aet) of 2 and 0.5% for

four heats of 2 1/4 Cr-1 Mo steel.

strain range was sensitive to heat treatment as well as to heat-to-heat variations (fig. 16). The data for three heats were generated at ORNL [ 11 ] on isothermal annealed steel; the Ellis et al. [31] data (heat 3) for a heat of steel in the annealed and isothermal annealed condition are also given. Note that one heat (50557) was a low carbon heat. Comparisons of the stress ranges obtained for a given strain range and temperature (fig. 16) reveal considerable heat-to-heat variability. Moreover, the strength of the response of a given heat compared with the others, remains consistent over the entire temperature range studied. Maximum differences in cyclic stress response between the various heats ococcurred over the temperature range of about 316°C (600°F) to 427°C (800°F). This is the temperature range where the influence of dynamic strain aging (fig. 10) is most pronounced. Comparing the cyclic stress response of heat 3 in

both the annealed and isothermal annealed condition, it is apparent that material in the isothermal annealed condition at both strain ranges required higher stress levels to achieve the indicated strain ranges than material in the annealed condition. Again differences tended to be most pronounced at test temperatures approaching the strain-aging temperature range. The isothermal annealed steel also showed higher monotonic yield and ultimate tensile strength values [31 ]. These results again show the need to understand the heat treatment effects from a broader view point than simply knowing whether the material is annealed or isothermally annealed. As a result of the observations in fig. 16 on the differences between the behavior of heat 3 in the annealed and isothermal annealed condition, Ellis et al. [31 ] concluded that the tensile properties of the isothermally annealed condition are superior to those of the materials in the annealed condition and the creep-fatigue properties of annealed

R.L. Klueh et al. / Elevated temperature reactor design

material are marginally superior. Obviously, the unqualified statement that the "tensile properties of 2 1/4 Cr-1 Mo steel in the isothermally annealed condition are superior to those of the material in the annealed condition" [31] is not true, for as seen above for heat 20017 (fig. 10), just the opposite was true. To understand this apparent descrepancy, one need only examine the two heat treatments [31]. The isothermally annealed steel was quickly cooled from 900 ° to 704°C, held for 2 h, then air cooled. The steel was annealed by a furnace cool from 900 ° to 704°C at about 28°C/h. It was held at 704°C for 4 h (actually, this also appears to be an isothermal anneal). Hence, the annealed steel transformed over almost an 11-h period, whereas the isothermally annealed steel transformed over a period of about 4 h, with most of the austenite to ferrite transformation occurring in this latter case at 704°C. Since the longer time at higher temperatures allows for more diffusion, less molybdenum and carbon are expected to remain in solution in the annealed steel. Thus, one should expect less dynamic strain aging. Furthermore, before the isothermal anneal treatment, the steel was cold rolled. This resulted in a small grain size, which also helps account for its greater strength [31 ] (the cold work could also affect the CCT diagram). Because of different isothermal anneal treatments, the cyclic strength properties for heat 3 also exceed those of the other three heats, which were heat treated as indicated by the curve for IA in fig. 8 (i.e., heat 3 spent less time at the elevated temperatures). As expected, carbon also plays a role in the dynamic strain-aging temperature regime. This is seen (fig. 16) by comparing the curve for the low carbon heat (50557) with the curves for the other three heats. There has also been a limited amount of work on the effect of heat-to-heat variation on the fatigue and creep-fatigue properties of type 304 stainless steel. In one study [32], when five different heats of type 304 stainless steel were subjected to continuous fully reversed strain-controlled cycling at 593°C and at a strain rate of 4 X 10-a/s, little difference in cyclic life times attributable to heat-to-heat variations was found. Similar results were obtained in another study on three additional heats [33]. Likewise, little difference could be seen in the cyclic strain-controlled tensile hold time behavior of four out of the five heats used in the first study [32]. However, the fifth

241

heat showed a remarkable resistance to the deleterious effects of tensile hold times in both the aged and annealed conditions, with an accompanying increased resistance to intergranular crack propagation. There is some indication that this latter result was due to a higher niobium content for that heat. To date, not enough long-time elevated-temperature creep-fatigue data have been generated to accurately define the heat-to-heat variations in the austenitic stainless steels. Such studies are complicated by the known influence of environment that occurs at elevated temperatures. Such environmental effects tend to mask heat-to-heat effects. Further, it is known that depending on conditions, metallurgical variations produced by pretest aging can either increase or decrease the fatigue life of material cycled under creepfatigue conditions. Aging produces carbide precipitates along grain boundaries; these precipitates inhibit intergranular crack propagation. Since variations in the chemical composition (e.g., the amount of carbon, nitrogen, and niobium present) can influence the precipitation of these carbides, it is expected that such differences would also influence cyclic lifetimes of components subjected to prolonged creep-fatigue interaction conditions. It appears, therefore, that more long-term tests accompanied by microstructural analysis are required to understand the extent to which heat-to-heat variations influence creep-fatigue properties.

5. Summary and conclusions If the designers of nuclear reactors and reactor components are to optimize their studies, they must be cognizant of the metallurgy of the alloys that are to be used in that design. Metallurgical effects that occur within an alloy during processing, fabrication, and service can give rise to changes in the mechanical properties Of the alloy. The types of effects that occur depend strongly on the type of alloy (e.g., precipitation processes affect the properties of 2 1/4 Cr-1 Mo steel to a much greater extent than they affect types 304 and 316 stainless steels). An understanding of the metallurgical causes for mechanical properties data variation and properties changes should eventually lead to design data that will more accurately represent the performance characteristics of the material

242

R.L. Klueh et al. / Elevated temperature reactor design

to be employed for a given reactor or reactor component. In this paper some of the types of mechanical properties variations observed for 2 1/4 Cr-1 Mo steel and types 304 and 316 stainless steels were examined. For the austenitic stainless steels, much of the effect can be traced to chemical differences (primarily carbon and nitrogen) and grain size. However, for 2 1/4 Cr-1 Mo steel much of the observed variation can be traced to the microstructures developed during the heat treatment and the evolution that the microstructure goes through during elevated-temperature exposure (during test or in service). Although most of this discussion has been directed toward tensile and creep properties with only a brief discussion of fatigue, explanations that involve these same metallurgical processes should also apply to observations on such other mechanical properties as relaxation, creep-fatigue interactions and subcritical crack growth.

Acknowledgements We gratefully acknowledge the review of the manuscript by K.C. Liu, J.M. Leitnaker, W.R. Martin, and D.A. Canonico; George Griffith for editing; and Kathryn Witherspoon for typing the manuscript.

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