Materials Science Forum Vols. 715-716 (2012) pp 813-818 Online available since 2012/Apr/12 at www.scientific.net © (2012) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/MSF.715-716.813
Microstructural evolution in 9%Cr heat resistant steels under creep conditions Rustam Kaibyshev Belgorod State University, Pobedy 85, Belgorod 308015, Russia email
[email protected] Keywords: creep, 9%Cr steel, carbides, Laves phase, martensite, subgrains, dynamic recovery.
Abstract. Microstructural design of a new generation of 9%Cr steels for fossil power plants is considered. It was shown that microstructural stability of 9%Cr steels impairs their creep resistance. Two types of restoration processes can occur in the heat resistance steels under creep conditions: (i) normal grain growth and (ii) dynamic recovery. The first process associates with the migration of high-angle boundaries (HAGB) of blocks of tempered martensite lath structure (TMLS). However, their migration is negligible even during creep deformation. Boundaries of packets and prior austenite boundaries (PAB) are effectively pinned by precipitations of M23C6 and Laves phase Fe2(W,N). The second process consists of transformation of lath boundaries to subboundaries and their subsequent migration (subgrain coarsening) under creep. Under aging the migration of lowangle boundaries (LAGB) is retarded by uniformly distributed nanoscale M(C,N) dispersoids and particles of M23C6 precipitated on these boundaries under tempering. Under creep the dissolution of M23C6 carbides located along LAGBs and coagulation of uniformly distributed M(C,N) carbonitrides facilitates LAGB migration. It was shown that the normal grain growth is not important for deterioration of creep strength. Conversion of the lath boundaries into subgrain boundaries strongly decreases creep rate. In contrast, continuous subgrain coarsening is the main process restricting the ability of the 9%Cr steel for long-range service under creep conditions. Tertiary creep is attained due to the occurrence of subgrain coarsening. Introduction Tempered martensitic 9%Cr steels are advanced materials for a new generation of fossil plants [1]. These steels gain their creep resistance mainly from a tempered martensitic lath structure (TMLS) that is essentially stable under service condition [2]. Excellent high-temperature strength of these steels is attributed to dispersion hardening and the strengthening due to TMLS. The last strengthening associates with the boundary hardening in accordance with the well-known HallPetch equation and dislocation hardening. Superposition of these three strengthening mechanisms provides superior creep resistance of the martensitic steels in comparison with traditional bainitic steels [1]. It was shown by numerous authors that the stability of TMLS is extremely important for creep resistance of the 9%Cr steels [1]; these steels are resistant to creep under exploitation conditions until TMLS remains essentially unchanged. Therefore, any improvement of the stability of TMLS under creep condition leads to enhancement of working temperature of the martensitic steels that is a key parameter for efficiency of fossil power plants. The purpose of the present work is to summarize reported experimental data and make clear the structural aspects of high creep resistance of the 9%Cr martensitic steels. In addition, the controlling factors for the stability of TMLS under creep conditions are considered. Specific attention is paid to contribution of dispersoids of different origin and solute elements in this stability.
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Recrystallization and Grain Growth IV
The lath martensitic structure Quantitative studies of TMLS of these steels using electron back scattering diffraction and transmission electron microscopy techniques showed that the TMLS of a 0.2% carbon steel [3] and the 9%Cr steels [4-8] is essentially the same. Difference in TMLS of these two types of steels resides in structural parameters. In general TMLS has a four-level hierarchy in its morphology, i.e. prior austenite grains (PAG), packets, blocks and laths (Fig.1) [3,4]. The martensite lath is a band delimited by low-angle boundaries (LAGB) with a high density of lattice dislocations (Fig.1). Several laths with the same crystallographic orientations aggregate to a block bounded by moderateto-high angle boundaries (Fig.1); the misorientation of block boundaries has to be larger than 10.53o [3]. Therefore, usually, blocks are bounded partly by high-angle boundaries (HAGBs) and partly by moderate-to-high angle boundaries with misorientation ranging from 10.53 to 15o. The packet is an aggregate of the blocks with the same plane. All boundaries of prior austenite grains (PAG) and packets and the majority of block boundaries are HAGBs in the 9%Cr steels [5,6]. It is worth noting that the portion of coincident site lattice (CSL) boundaries is relatively low and does not exceed 10pct [5,6]. A prior austenite grain composes several packets (Fig.1).
Fig.1 Schematic presentation of the lath martensite structure The size of PAGs usually varies from 10 to 16 µm in the 9%Cr steels [8,9]. Dimensions of packets and blocks could not be determined with high accuracy due to the fact that a limited number of areas could be analyzed by careful orientation analysis [3]. Therefore, it is reasonable to use the term “grain” to describe a microstructural unit which is entirely delimited by HAGBs; grain size for TMLS means a distance between the two nearest HAGBs. Average size of these grains is ~2µm and, therefore, the 9%Cr steels are ultra-fine grained materials. It is quite reasonable to distinguish a difference between the terms of PAG and the grains. It is caused by a difference in size of particles of M23C6 carbides and Laves phase Fe2(W,Mo) precipitated on prior austenite boundaries (PAB) and the grain boundaries. Under creep condition the size of PAG remains unchanged, while the grains tend to grow with a low rate [5,6,8,9]. The other important structural parameters of TMLS on micro-scale level are thickness of martensite laths which varies from 200 to 500 nm, density of lattice dislocations, ρ, and internal elastic strain, ∆a/a [5-9]. Usually, an average thickness of ~350 nm [5-8] is observed in TMLS. Low values of the thickness were found [7] in the quenched martensitic structure, while the highest values of thickness were observed after long-range aging under creep conditions [6,8]. Notably the martensite laths gradually transform to subgrains having essentially equiaxed shape during creep [2,5,6,8,9]. An average subgrain size is always higher than the prior thickness of martensitic laths [2,5,6,8,9]. Average lattice dislocation density in TMLS is ~4×1014 m-2 [5,6,8,9]; ∆a/a is ~0.3 after
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quenching and ~0.1 after standard tempering at 750-780oC for 3 hours [7]. Average density of lattice dislocations remains almost unchanged under long-range aging [5,6,8]. In contrast, the final density of lattice dislocation after rupture may be 10 times lower than that after the long-range aging at the same temperature [5,6,8-10]. There exist no data on effect of creep deformation or/and longrange aging on the internal elastic strain, ∆a/a. It is worth noting that some authors [11] consider the thickness of martensitic laths in terms of a subgrain size. However, from our point of view it is not correct. The lath can be considered as an element of low-energy dislocation structure (LEDS) (Fig.2a) [12] due to the fact that the lath boundary is an irregular array of dislocations. Stored energy in the lath boundary is higher than that in a subboundary with a similar misorientation by a factor of about 5. Therefore, there exists a high driving force for migration of lath boundaries. Thus, 9%Cr martensitic steels are materials with high stored energy in forms of high dislocation density and numerous high energy grain boundaries. The last term includes the energy of lath boundaries and HAGBs of blocks, packets and PAGs. Therefore, there exists a potential associated with high driving forces [12] for occurrence of normal grain growth or/and recovery under creep conditions. Obviously, the occurrence of the recrystallization process in heat-resistance 9%Cr martensitic steels leads to the deterioration of their creep strength [1]. Alloying philosophy of these steels is aimed to suppress recrystallization and recovery processes by precipitations of carbides, carbonitrides, Laves phase and to obtain low diffusivity and solution hardening by increased solute content [1]. The 9%Cr martensitic steels are a unique material in which the ultra-fine grained structure, well-defined subgrain structure and high dislocation density are prevented from recovery and recrystallization at intermediate temperatures under creep conditions by high volume fraction of nano and submicrometer scale dispersoids and solutes having high ability to decrease diffusivity. It seems [5-9] that migration of HAGBs is effectively hindered by second phase particles and solutes; normal grain growth plays an unimportant role in microstructural evolution of TMLS. Creep behavior is controlled by occurrence of dynamic recovery which actually consists of two stages: (i) transformation of the lath boundaries into subgrain boundaries (Fig.2b) and (ii) continuous subgrain coarsening (Fig.2c). The exact mechanism of the last process is unknown. Role of dispersoids and solutes in microstructural evolution TMLS. Phase distribution of dispersoids in TMLS is schematically presented in Figs.2a&3a [8]. The fine M(C,N) carbonitrides are uniformly distributed within the ferritic matrix. Two types of these carbonitrides can be distinguished [8]: Nb-rich M(C,N) particles having an equiaxed shape and size ranging from 25 to 40 nm and V-rich M(C,N) dispersoids having plate-like shape with a thickness of ~8 nm. The volume fraction of Nb-rich carbonitrides is ~0.2% that is significantly higher than the equilibrium fraction of these dispersoids calculated by ThermoCalc. In contrast, the volume fraction of V-rich carbonitrides is
