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ISSN 0031918X, The Physics of Metals and Metallography, 2010, Vol. 109, No. 2, pp. 186–200. © Pleiades Publishing, Ltd., 2010. Original Russian Text © R.O. Kaybyshev, V.N. Skorobogatykh, I.A. Shchenkova, 2010, published in Fizika Metallov i Metallovedenie, 2010, Vol. 109, No. 2, pp. 200–213.

STRENGTH AND PLASTICITY

New Martensitic Steels for Fossil Power Plant: Creep Resistance R. O. Kaybysheva, V. N. Skorobogatykhb, and I. A. Shchenkovab a

Belgorod State University, ul. Pobedy 85, Belgorod, 308015, Russia Central Research Institute of the Machine Building Technology, ul. Sharikopodshipnikovskaya 4, Moscow,115088 Russia

b

Received December 2, 2008

Abstract—In this paper, we consider the origin of hightemperature strength of heatresistant steels belong ing to martensitic class developed on the basis of the Fe–9%Cr alloy for the boiler pipes and steam pipelines of power plants at steam temperatures of up to 620°C and pressures to 300 atm. In addition, we give a brief information on the physical processes that determine the creep strength and consider the alloying philosophy of traditional heatresistant steels. The effect of the chemical and phase composition of heatresistant steels and their structure on creep strength is analyzed in detail. It is shown that the combination of the solidsolu tion alloying by elements such as W and Mo, as well as the introduction of carbides of the MX type into the matrix with the formation of a dislocation structure of tempered martensite, ensures a significant increase in creep resistance. The steels of the martensitic class withstand creep until an extensive polygonization starts in the dislocation structure of the tempered martensite(“troostomartensite”), which is suppressed by V(C,N) and Nb(C,N) dispersoids. Correspondingly, the service life of these steels is determined by the time during which the dispersed nanocarbonitrides withstand coalescence, while tungsten and molybdenum remain in the solid solution. The precipitation of the Laves phases Fe2(W,Mo) and the coalescence of carbides lead to the development of migration of lowangle boundaries, and the steel loses its ability to resist creep. Key words: martensitic steel, hightemperature strength, carbides, Laves phases DOI: 10.1134/S0031918X10020110

INTRODUCTION At present, in the United States, Japan, China, and European countries coalfired power plants with ultrasupercritical (USC) steam parameters (Т = 600– 620°С, Р = 250–340 atm) [1–5] are intensely built. The transition to such steam parameters from the tra ditional ones (Т = 545°C, Р = 240 atm) makes it pos sible to increase the efficiency of coalfired power plants from ~35% to ~44%. This transition became possible after the development in the United States, Japan, and Europe of new heatresistant martensitic steels on the basis of the Fe–9% Cr steel as a result of intensive R&D works that were carried out for the last 25 years. These steels have not only high creep resis tance, but also a relatively low cost, which ensures the high economic efficiency of their use. The other important merit is a low thermalexpansion coeffi cient (TEC), which ensures the high cyclic load capa bility of the power units in which these steels are used. In the modern boiler constructions, they are employed as the material for the hightemperature circuits of superheated steam, collectors, main steam pipelines, steam turbines. These steels are joined with the classi cal heatresistant steels of the P22/23/24 types by composite welds, allowing making heavier parts of the boilers from pearlitic or bainitic steels, which is of high economic importance.

At present, a whole number of largescale research projects devoted to the development and implementa tion of martensitic steels have been initiated in Rus sian Federation [4]. The machinebuilding and metal lurgical plants of RF begin to develop the production of these steels and fabrication of parts from them. These works are retarded by the lack of technical information. The aim of this review is to briefly con sider the latest achievements in the metallurgical background of martensitic heatresistant steels in the Russianlanguage literature and to describe the basic ways for improving the hightemperature properties of these steels. This review is restricted to the consider ation of boiler steels. We also do not analyze some steels with an increased content of Cr of the P122 (Sumitomo) type, since they exhibit only unessential differences from the steels of the P92 (Nippon Steel) and P91 types. CREEP OF MATERIALS: GENERAL CONCEPTS The basic characteristic of hightemperature steels employed in fossil power plants is their resistance to creep, which is characterized by two quantities: creep strength, i.e., a constant stress that causes (in a specific time at an assigned temperature) strain that does not exceed an assigned magnitude (for example, 1% for 105 h at 600°С); and rupture time.

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The creep strength is determined by the strain rate at the second (steadystate) stage. The less the creep rate and the stronger its functional dependence on the applied stress, the higher the creep strength. The strain rate obeys two different laws [6–10]. At low tempera tures, the deformation behavior of a material is described by the equation – Q⎞ , ε· = B exp ( βσ ) exp ⎛  ⎝ RT⎠

(1)

where ε· is the strain rate, В and β are constants, σ is the stress at the steadystate stage, Q is the activation energy for deformation, R is the universal gas con stant, and T is the absolute temperature [6–11]. In the hightemperature range, the strain rate is described by a power law n – Q⎞ ε· = Aσ exp ⎛  , ⎝ RT⎠

(2)

where A is a constant, and n is the stress exponent. The magnitude of n is equal to 1 in the case of viscous creep; to 2 in the regime of superplasticity; and varies from 4 to 5 in the region of hot deformation and from 6 to 8 in the region of warm deformation. The transi tion from hot to warm deformation leads to an increase in the magnitude of the exponent according to the rule n = n + 2 [6–10]. Upon the transition from the warm to cold deformation, the deformation behavior of material ceases to be described by Eq. (2); the dependence of the strain rate on the applied stress begins to obey Eq. (1). Note that the determining rela tionships (1) and (2) with an identical accuracy describe the dependence of the strain rate on the applied stress at n = 7–9 [12]. Consequently, the value n ~ 8 in Eq. (2) should be considered as a minimum value at which the material can be employed as high temperature. For convenience of the graphic repre sentation, the dependence of the strain rate on the applied stress at the steady stage of creep is plotted in logarithmic coordinates, namely, as a normalized strain rate ε· kT(D1Gb) (where k is the Boltzmann con stant, D1 is the temperaturedependent coefficient of lattice diffusion, G is the temperaturedependent shear modulus, and b is the Burgers vector of disloca tions) versus the normalized stress σ/G (Fig. 1) [6–11]; if the steadystate stage is absent, then the strain rate ε· ss is defined as the minimum strain rate upon a pas sive loading or the maximum flow stress σpeak upon active load. Figure 2 displays the temperature depen dence of the activation energy for plastic deformation Q. For the majority of materials in which the creep rate is lower than the slip rate, the magnitude of Q in the region of cold deformation linearly grows with tem perature until it becomes equal to the activation energy for pipe diffusion. In the region of warm defor mation, Q is equal to the activation energy for pipe dif fusion, while in the region of hot deformation, it THE PHYSICS OF METALS AND METALLOGRAPHY

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becomes equal to the activation energy for lattice dif fusion. According to the modern concepts [11–14], the changes in the activation energy and in the determin ing relationships of creep (1) and (2) with a change in temperature or applied stresses are connected with changes in the mechanisms of plastic deformation, which for the coarsegrained (≥10 μm) materials can be briefly presented as follows (Fig. 1 and 2). (1) In the region of low stresses, creep occurs due to the diffusion of vacancies through the volume of grains. The activation energy is equal to the activation energy for lattice diffusion (Nabarro–Herring, or Harper–Dorn creep). The development of this mech anism of creep in the process of exploitation of the hightemperature materials is inadmissible. (2) In the region of hot deformation, the disloca tions can rearrange in directions perpendicular to the plane of their slip through large distances due to the hightemperature climb, which is controlled by lattice diffusion. This gives the possibility for dislocations to bypass obstacles and leads to low values of creep strength. The regime of hightemperature materials exploitation cannot coincide with the strainrate– temperature interval of hot deformation. (3) In the region of warm deformation, the disloca tions rearrange to only small distances due to the slow lowtemperature climb controlled by pipe diffusion. Under specific conditions, which will be described below, the regime of exploitation of superalloys can coincide with the strainrate–temperature interval of warm deformation. (4) In the region of cold deformation, the disloca tions have no possibility to rearrange; during their motion, they are forced to intersect obstacles. The tem peratures and stresses corresponding to the operating conditions for the hightemperature materials must be located within the interval of cold deformation. From the viewpoint of the physical mechanisms of deformation, an increase in the temperature of the operation of hightemperature materials is achieved by shifting the boundaries between different tempera ture–strainrate regimes of deformation (Fig. 1) toward higher temperatures. The most effective way to achieve this is the solidsolution alloying with ele ments that increase the binding force between the atoms and, correspondingly, decrease the diffusion rate. Thus, the alloying of aluminum with 6% Cu leads to an increase in the temperature boundary between the warm and hot deformation by 200 K [14]. This means that the solidsolution alloying is an efficient method of increasing maximum temperature of high temperature material operation. Another method of increasing the operating temper ature is the introduction of dispersed nanoparticles resistant to the coalescence into the material; this leads to the appearance of threshold stresses below which no deformation occurs (Fig. 3) [11–14, 15–17]. In this Vol. 109

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10–4 10–5 Cold deformation

10–6

εkT/(D1Gb)

10–7

Sherby–Burke criterion

10–8 n=n+2 10

Warm deformation

–9

n=5

10–10

Hot deformation

–11

10

10–12

n=1

10–13

Diffusion creep

10–14

10–4

10–3 (σ – σth)/G

10–2

Fig. 1. Dependence of the normalized strain rate ε· kT/(D1Gb) on the normalized stresses (σ − σth)/G [6] (k is the Boltzmann constant; D1 is the temperaturedependent coefficient of latlice diffusion; G is the temperaturedependent shear modulus; and b is the Burgers vector).

case, the deformation behavior of a material is described by the following equation: n ε· = A [ ( σ – σ th )/G ] exp ( – Q c /RT ),

(3)

where σth is the threshold stress. It is assumed [18] that the magnitude of the threshold stresses is approxi mately equal to the applied stresses at which ε· = 5 × 10–8 s–1, and the exponent in Eq. (2) is n > 20 (in Eq. (3), n has the “true” values that are obtained by specific calculations [11–14, 15–17] and lie within the limits of n = 4–8). This means that the dispersed particles, which are the “sources” of threshold

stresses, suppress lowrate creep. The use of dispersion strengthening makes it possible to employ the temper ature–strainrate operating conditions of hightem perature materials that coincide with the interval of warm deformation. Dislocations cannot rapidly go around particles with dimensions of 10–40 nm by climbing and are forced to overcome them by generat ing loops, which requires additional stresses, since an increase in the dislocation length means an increase in its energy. The threshold stresses are inversely propor tional to the distance between the particles [11–14, 15–17]. An increase in the fraction of the dispersed particles of the second phases and/or their refinement

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Q

ε, s–1

Qbd Qpd

189

Cold deformation

Warm deformation

.

ε ~ σn

Hot deformation 10–6

Material with dispersed particles

Т 5 × 10–8 Fig. 2. Fig, 2. Temperature dependence of the activation energy for plastic deformation: Qbd, activation energy for lattice diffusion; and Qpd, activation f energy or pipe diffu sion [6].

leads to an increase in the threshold stresses. If we select the volume, the distribution, and the average size of particles in such a way that the threshold stresses will be greater than the applied stresses, then it is possible to completely suppress the movement of dislocations. It should, however, be noted that the threshold stresses suppress only dislocation creep. The development of diffusion creep leads to the disappear ance of threshold stresses [19]. Thus, for a hightemperature material to have a maximum creep strength, it must have a matrix struc ture consisting of a heavily alloyed solid solution with distributed dispersed nanosized particles. An increase in the degree of the alloying of solid solution with ele ments that reduce the diffusion rate, just as an increase in the fraction and a decrease in the size of nanoparti cles must lead to an increase in the operating temper ature of the hightemperature materials. The grain size quite significantly affects the heat resistance of steels and alloys; this is connected with its strong effect on the rate of diffusion creep, which occurs at low strain rates. The rate of the most com mon (Nabarro–Herring) mechanism of diffusion creep is expressed by the following equation [6–8]: D 1 σΩ⎞  , ε· NH = B NH ⎛  (4) ⎝ d 2 kT ⎠ where Ω = 0.7b3 is the atomic volume, d is the average grain size, and BNH = 15 is a dimensional constant. The rate of diffusion creep is inversely proportional to d2. Consequently, an increase in the grain size effec tively reduces the rate of diffusion creep. Correspond ingly, the condition necessary for high creep strength, which can be obtained only in the regime of disloca tion creep, is a large grain size, which makes it possible to suppress diffusion creep or to move its interval toward smaller applied stresses. The rupture time τ is the second most important characteristic of the creep behavior. In the hightem THE PHYSICS OF METALS AND METALLOGRAPHY

Solid solution

σth

σ

Fig. 3. Influence of dispersed particles on the dependence of the strain rate on the applied stresses [6, 15, 16].

perature alloys, in which at grain boundaries there are no secondphase precipitates with a high (in compar ison with the matrix) shear modulus, this time is inversely proportional to the creep rate at the steady stage (to the minimum creep rate) according to the Monkman–Grant relationship t = C/ ε· ss , (5) where С is a constant. The quantity τ correlates with the relative elongation during creep. These character istics have a fundamental importance for the service properties of hightemperature materials. It is pre cisely the magnitude of τ that is used for estimating the life assessment and the allowable stresses for heat resistant steels [20–21]. The rupture time is determined by many factors [6–8, 20]; however, the main of these is the formation of straininduced pores at grain boundaries. This is related to the operation of grainboundary sliding (GBS) [6, 22]. The rotation of grains relative to each other leads to the formation of pores if there is no GBS accommodation, which is possible due to the diffusion or/and slip and climb of lattice or grainboundary dis locations [6, 22]. The possibility for accommodation during GBS is determined by the grain size and by the relationship between the critical stresses for slip and GBS (τs/τgbs) [20]. The small grain size and the low relationship provide easier accommodation of GBS by dislocation slip [22]. There are known several mechanisms of pores for mation at grain boundaries. All of these are connected with the existence of structural defects at grain bound aries, such as the facets, which arise as a result of the intersection of a grain boundary with slip bands (Fig. 4a), triple junctions, intergranular particles with a shear modulus greater than that of the metallic matrix (Fig. 4b), etc. [6, 22]. These defects lead to the appearance of local stresses upon the displacement of grains relative to each other as a result of GBS that Vol. 109

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(a) σ

(b)

τgbs σn pore τgbs

Secondary slip Cr23C6 or Fe2W GBS GBS

pore Primary slip

Fig. 4. Mechanisms of pore formation during creep [6, 7, 22]: (a) the formation of pores at grainboundary steps, at the sites of emergence of slip bands; and (b) the formation of a pore at a particle located at a grain boundary.

cannot be rapidly relaxed, which leads to the appear ance of pores remaining stable during deformation. It is worth noting [20] that the probability of the origin of pores at large hard particles located along grain boundaries (Fig. 4b) is almost 100 times higher than the probability of their nucleation at usual grain boundaries (Fig. 4a). In fact, Eq. (5) is fulfilled if the pore formation occurs via the mechanism shown in Fig. 4a. When large undeformable particles are formed at grain boundaries, the failure occurs at times that are less than those predicted by Eq. (5). Correspondingly, the precipitation of large carbides of the M23C6 type at grain boundaries in heatresistant steels leads to a pre mature fracture of material. Thus, to increase the rupture time τ, the following conditions should be fulfilled. (1) GBS should be suppressed. This is achieved by increasing stresses of 105 by the microalloying of high temperature steels and alloys with boron, which forms equilibrium segregates at grain boundaries [19], or by the generation of precipitates of dispersed particles less than 10 nm in size at and/or near grain boundaries [23]. (2) The decrease in the grain size facilitates accom modation during GBS; i.e., the optimum grain size of hightemperature materials is a result of a compromise between the suppression of diffusion creep and the need of avoiding pore formation at grain boundaries. (3) It is necessary to avoid the precipitation of large solid particles at grain boundaries. If we avoid a premature failure in the process of creep as a result of formation and growth of strain induced pores, then the allowable creep stresses will be determined by the creep rate and, correspondingly, by

the solidsolution alloying and also by the volume fraction and by the size of dispersed particles. TRADITIONAL HEATRESISTANT STEELS During the last 50 years, the main materials for the boiler tubes and steam pipelines are steels of the pearl itic class, such as 12Kh1MF, 15Kh1MF, P22 (in the Englishlanguage literature these steels are referred to steels of bainitic class [24]) and bainitic steels P23 and P24 (Tables 1 and 2). Their chemical composition and the creep stress at which they retain operational capa bility for 105 h [8–12] are given in Tables 1 and 2. With all their differences in the chemical and phase compo sition, these steels have much in common from the viewpoint of the alloying philosophy, the route of heat treatment, and the resultant structure. The route of heat treatment consists of normalization with a subse quent hightemperature annealing. The normaliza tion for these steels in the case of thinwall pipes leads to the formation of bainite or even martensite, i.e., in fact, quenching. From the viewpoint of the theory of heat treatment, the final operation for these steels is hightemperature tempering, during which there occurs both a complete decomposition of martensite or of bainite, and a recrystallization. These steels after the heat treatment before the beginning of exploita tion have an equilibrium structure, which consists of polygonal ferrite grains and precipitates of carbides inside them (Fig. 5a). It should be noted that this leads to an interesting phenomenon: the regime of heat treatment affects the short term characteristics of creep and does not influence the rupture time and the

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Table 1. Chemical composition of traditional heatresistant steels Steel

C, %

Cr, %

W, %

Mo, %

P22 12Kh1MF 15Kh1MF P23 P24

0.05–0.15 0.08–0.15 0.1–0.16 0.04–0.1 0.05–0.1

1.9–2.6 0.9–1.2 1.1–1.4 1.9–2.6 2.2–2.6

– – – 1.45 –

0.87–1.13 0.25–0.35 0.9–1.1 0.05–0.3 0.9–1.1

Steel

V, %

Nb, %

Si, %

Mn, %

B, %

P22 12Kh1MF 15Kh1MF P23 P24

– 0.15–0.3 0.2–0.25 0.2–0.3 0.2–0.3

– – – 0.02–0.08 –

0.5

0.3–0.6



1%) content of W holds in control the coalescence of carbides of the М23C6 type and slows down the precipitation of the Laves phases of the (Fe, Cr)2(W, Mo) type during creep. (3) The introduction of boron in concentrations of more than 0.003% ensures its segregation in the M23C6 carbides located at grain boundaries. This substantially increases the resistance of these carbides to coales cence, which increases rupture time during creep. (4) The introduction of V and Nb into steel leads to the precipitation of nanosized carbonitrides of the MX type. They suppress the migration of lowangle and highangle grain boundaries and the movement of dis locations during creep. Simultaneously, the Nb(C,N) dispersoids make the steel hereditarily finegrained. ACKNOWLEDGMENTS The authors thank the Federal Agency for Science and Innovations for the support of this work within the framework of the Federal Project no. 02.523.12.3019 (“Development of Nanostructured HighTempera ture Steels and Technologies of the Production of HighTemperature Elements of Power Equipment of New Generation from These Steels”).

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THE PHYSICS OF METALS AND METALLOGRAPHY

Vol. 109

No. 2

2010