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High-chromium martensitic steels are one of the basic creep-resisting construction materials used for the modernization of old and the construction of new power units. During the service under creep conditions, the metastable microstructure of martensitic steels undergoes gradual degradation. The rate of degradation mostly depends on the operating temperature, but it is also affected by stresses. The changes in the microstructure of martensitic steels have an influence on the decrease in their mechanical properties, including creep resistance. The knowledge and description of the changes in the microstructure of steels working under creep conditions allow extending the time of safe operation of the elements of power systems. The paper presents and describes the main mechanisms of degradation of 9–12%Cr martensitic steels on the basis of the independent studies and literature data.
Institute of Materials Engineering, Czestochowa University of Technology, Czestochowa, Poland
Cezary Kolan
Institute of Materials Engineering, Czestochowa University of Technology, Czestochowa, Poland
Joanna Jasak
Institute of Materials Engineering, Czestochowa University of Technology, Czestochowa, Poland
*Address all correspondence to: grisza@wip.pcz.pl
1. Introduction
The need to reduce emissions of pollutants (mainly CO2) to the atmosphere enforced by increasingly stringent EU directives has contributed to the development of conventional energy. Restrictions on emissions to the atmosphere caused by the combustion of fossil fuels have forced the power industry to increase the thermal efficiency of power units (from 33–35 to 40%, and ultimately up to 50%). On the other hand, the need to increase the thermal efficiency of power units involves a significant increase in steam parameters (pressure, temperature). This requires the construction of new and the modernization of existing power units to allow them to operate at the so-called supercritical or ultra-supercritical steam parameters. The increase in steam parameters in new and existing power units was possible due to the materials revolution in the power industry and was associated with the introduction of new grades of steels and cast steels with higher resistance to creep and oxidation than that of the materials used so far. The implementation of new materials in the power industry took place primarily through modifications and optimizations of steels that were already being used in the power sector. It has contributed not only to the increase in steam parameters, but also to the reduction in overall dimensions of boiler components, and thus in their weight, which has also a significant impact on the reduction in the energy production costs [1, 2].
One of the new steel groups introduced to the power industry was high-chromium martensitic steels containing 9–12%Cr. By the optimization of carbon content and the introduction of additions and micro-additions such as W, Co, V, Nb, N, B, and Cu to these steels, the construction materials characterized by high mechanical properties were obtained. For example, their creep strength is higher by approx. 20–25% than that of the steels used so far [2, 3]. The expected high reliability and long life of up to a minimum of 200,000 h of pressure parts made from, among others, 9–12%Cr steels require understanding and describing the effects and microstructure degradation processes for these materials. Based on many years of Authors’ own research and literature data, the main steel/martensitic cast steel microstructure degradation mechanisms and their impact on mechanical properties were described and characterized in this paper.
2. Microstructure degradation and properties of 9–12%Cr steels
The basic requirement for creep-resistant steels used in the broadly understood power industry is to maintain—for a relatively long time of operation (at present, 200–250,000 h)—the assumed mechanical properties at the operating temperature of power equipment components. The maintenance of the required mechanical properties of creep-resistant steels during long-term service depends on the stability of their microstructure. The structural components of the power equipment are influenced by the continuous destruction process, which has a significant impact on the life and time of safe service of a specific component. Therefore, the time of safe service for devices used in the power industry is one of the most important parameters related to their life, and it determines their applicability in this sector [1, 4, 5].
During their long-term service, progressive changes in microstructure of creep-resistant steels take place—the process of degradation of their microstructure occurs. For 9–12%Cr martensitic steels, the main microstructure degradation mechanisms include [6–8]:
matrix recovery and polygonization processes,
coagulation of M23C6 carbides,
precipitation of secondary phases: Laves phase and Z-phase, and
depletion of alloying elements in matrix.
Martensitic steels in the hardened condition are characterized by high dislocation density within the martensite laths (1016–1018 m−2). Due to high-temperature tempering of hardened steel, a more thermodynamically stable microstructure with still high dislocation density at 1012–1014 m2 (including free dislocations) within the subgrains formed during the tempering is obtained. The dislocation substructure in 9–12%Cr steels is characterized by small elongated subgrain (of 200–400 nm in width) and low-angle boundaries [6, 7, 9, 10]. High dislocation density and microstructure refinement with dislocation boundaries has a very intensive impact on the 9–12%Cr steel hardening with the dislocation hardening mechanism and the grain boundary hardening mechanism, respectively. The calculations in [11] showed that the gain in yield strength in martensitic steels for the above-mentioned mechanisms is 18 and 33%. In addition to this hardening, the following further mechanisms are additionally used to form the structure and mechanical properties of 9–12%Cr steels: solution hardening with interstitial and substitution elements and precipitation hardening [9, 11, 12]. In high-chromium martensitic steels, the precipitation hardening mechanism is mainly performed by the secondary particles precipitated when tempering M23C6 carbides and MX precipitates. In 9–12%Cr creep-resistant steels, three types of MX precipitates can occur [3, 13, 14]:
Secondary lamellar VN (VX) nitrides (carbonitrides), which are precipitated within the martensite laths during high-temperature tempering
Precipitate complexes consisting of the spherical NbX precipitate in which the VN precipitate nucleates, referred to as the “V-wings”
The degree of hardening with secondary phase precipitates depends mainly on the amount and size of precipitates and their distribution within the matrix.
High-chromium martensitic steels in the as-received condition (i.e., after quenching and tempering) have a metastable microstructure, which will be affected by gradual evolution as a result of long-term service (Figure 1).
Figure 1.
The microstructure of GP91 cast steel in the as-received condition [8].
Long-term effect of the temperature and time, in the case of creep as well as stress, leads to a decrease in strengthening with the dislocation mechanisms and the grain boundaries. A decrease in free dislocation density within grains and increase in size of the subgrains take place. The decrease in dislocation density with the time of service/aging is associated with the progressing process of their regrouping, arranging, annihilation, and entangling the dislocations in grain boundaries as well as the formation of cellular dislocation and subgrain microstructure (Figure 2). The matrix recovery and polygonization process results in the disappearance of martensite lath microstructure and the formation of polygonised ferrite microstructure Figures 3 and 4.
Figure 2.
The interaction of dislocations with lath/subgrain boundaries, PB2 steel, TEM.
Figure 3.
The microstructure of: (a) T91 steel after long-term service [12], (b) GP91 cast steel after 70,000 h aging at 600°C.
Figure 4.
The microstructure of polygonised ferrite with numerous precipitates in GP91 cast steel after low-cycle fatigue at 600°C with strain amplitude of 0.60%.
In the microstructure of martensitic steels, the formation of polygonal structure during service takes place due to the progressive increase in the size of subgrains. This process is slow because of low mobility of these subgrains. The stability of subgrain size has a positive impact on the maintenance of high mechanical properties, including creep resistance [16]. The increase in the size of subgrains occurs due to the migration or coalescence of the sub-boundaries.
The increase in the size of sub-boundaries usually takes place with the “Y” mechanism [15, 17, 18]. The migration with this mechanism is based on the movements of “Y” nodes that are the place where three sub-boundaries meet, which allows the coalescence of two low-angle boundaries. The increase in the size of subgrains with the “Y” node movement mechanism is shown in Figure 5.
Figure 5.
The morphology of “Y” nodes in: (a) P91, (b) PB2, TEM [15].
The matrix recovery and polygonization process takes place in the presence of secondary dispersion phases, which act as a stabilizing agent. The lath microstructure stability depends on the stability of M23C6 carbides precipitated on these phases (Figure 6). M23C6 carbides precipitate on the tempered martensite lath boundaries and on the subgrain boundaries preventing their growth due to the matrix polygonization, repolygonization, and recrystallization processes [17, 19]. In addition, the elongated shape of M23C6 carbides precipitated on the sub-boundaries (Figures 1, 3, and 5) has a positive impact on anchoring the boundaries by them as their contact surface with the boundary on the same volume fraction is bigger than for spherical particles [18].
Figure 6.
The interaction of dislocations in GP91 cast steel with particles of precipitates occurred after low-cycle fatigue.
The thermodynamic thermal stability of M23C6 carbides is not too high—the Cr23C6 carbide formation enthalpy is: –20 kJ/mol [20]. In the as-received condition, the size of M23C6 carbides in martensitic steels is 50–150 nm [8, 21]. The long-term service/aging contributes to changing the morphology of M23C6 carbides. These precipitates show a high tendency to coagulation. The process of coagulation of precipitates is determined by two basic factors: the thermodynamic and the kinetic one. The thermodynamic factor results from a large value of the surface energy of the interphase boundaries. As a result of coagulation, the surface energy decreases and aims at reaching the energy equilibrium. The kinetic factor of coagulation, on the other hand, is connected with the diffusion and reactions occurring on the boundary surface. They run at different rates, and the slowest one determines the rate of particle growth in the system, and thereby determines the kinetics of coagulation. The constant of the rate of growth of particles Kp in the matrix of martensitic steels is presented in Table 1.
Precipitate
Steel
Temperature, °C
Based on solubilities at tempering temperature
Based on solubilities at exposure temperature
Kp, m3 s−1
γ, J/m2
Kp, m3 s−1
γ, J/m2
MX
P92
600
1.17 × 10−32
0.5
8.58 × 10−33
0.5
650
9.5 × 10−32
0.5
65.5 × 10−33
0.5
Laves phase
P92
600
–
–
2.91 × 10−31
1.0
650
41.6 × 10−31
1.0
M23C6
P92
600
0.12 × 10−29
0.1
1.88 × 10−30
0.1
650
1.37 × 10−29
0.1
4.78 × 10−30
0.1
P91
600
2.88 × 10−29
1.0
7.67 × 10−30
0.5
650
25.3 × 10−29
0.8
59.8 × 10−30
0.3
Table 1.
Calculated coarsening rate constants Kp of MX, Laves phase, and M23C6 precipitated in P91 and P92 steel based on the shown fit values for the interfacial energy γ [22].
The coagulation of M23C6 carbides reduces their amount with almost the same volume fraction and results in the increase in distance between these precipitates. Also, nonuniform distribution of these precipitates within matrix makes them become a less effective factor controlling the increase in the size of subgrains (Figure 7). The literature data [23] show that the subgrain boundaries with mutual disorientation angle of less than 20° are not the points of preferential carbide precipitation. According to the research in [24], only about 8% of M23C6 carbides was precipitated at the low-angle boundaries with mutual disorientation angle of 8–15°. The low-angle boundaries represent at least 60% of the total amount of boundaries in tempered martensite. This limits their role and reduces their effectiveness as a substructure stabilizer, which results in the reduction in creep strength. The precipitate-free boundaries show higher mobility, which results in an increase in their width. The increase in the size of M23C6 carbides precipitated at the boundaries is also conducive to the reduction in ductility of 9–12%Cr steels [6, 8, 22–26].
Figure 7.
The microstructure of martensitic steels with both wide and narrow martensite laths visible: (a) P91, (b) PB2, TEM.
In steel with micro-addition of boron, the carbon atoms in M23C6 carbides are partially replaced by boron during the tempering, which results in the formation of M23(C, B)6 carboborides. Like M23C6 carbides, these precipitates occur at the grain boundaries and at the martensite lath boundaries. However, these precipitates are more finely dispersed and characterized by higher thermodynamic stability compared to M23C6 carbides [9, 14, 27]. This results in a slower increase in the size of these precipitates, which has a positive effect on the stability of tempered martensite lath microstructure and results in a higher creep resistance. Also according to [24], vanadium plays a significant role as a factor controlling the process of coagulation of M23C6 carbides. Vanadium dissolved in the matrix is conducive to decreasing of the coefficient of chromium diffusion in ferrite. Similar influence is also observed in the case of tantalum [28]. The temperature of work has a considerable effect on the rate of coagulation of M23C6 carbides. Elevating the temperature of service by 50°C can cause a growth of the rate of coagulation of these precipitates even by an order of magnitude.
The martensitic steels gain high creep resistance mainly due to the precipitation hardening provided by: MX nitrides, carbonitrides (where: M = V, Nb; X = C, N). MX precipitates are characterized by nanometric dimensions of about 10–50 nm, and in spite of their low volume fraction of 0.020–0.025, they ensure very strong hardening of creep-resistant steels (Figure 8).
Figure 8.
The MX precipitates in T91 steel after service: (a) NbC and VX, (b) V-wings [12].
The hardening with these precipitates is ensured by anchoring and hindering the motion of dislocations [6, 7, 9, 14, 15, 18, 29]. The calculations made for the P91 steel showed that the stress required for dislocation to “bypass” the carbide and nitride particles with the Orowan mechanism is as follows: for M23C6—39 MPa, for NbC—15 MPa, and for VN—106 MPa [30]. The MX precipitates in 9%Cr steels have a very high thermal stability (Table 1). The approximate formation enthalpy for these precipitates is as follows: for VC and NbC carbides, 55 and 70 kJ/mol, respectively, and for VN nitride, 125 kJ/mol [19]. High stability of MX precipitates and their coherent (semi-coherent) interphase boundaries cause that after approx. 100,000 h creep at 600°C, their size is similar to that as in the as-received condition [24, 30, 31].
In martensitic steels used at above 600°C and containing at least 10%Cr, MX precipitates represent a metastable phase and undergo a transition into a more thermodynamically stable Z-phase (Figure 9a) [30–34]. The disappearance of finely dispersed MX precipitates in the microstructure of these steels due to the MX carbide → Z-phase transition (complex Cr(V, Nb)N nitride) results in a very fast decrease in creep resistance [31, 32, 34, 35].
Figure 9.
Z-phase precipitate in T91 steel after service (a), interaction of dislocations inside the subgrain with MX precipitates (b).
According to [32, 35], one large Z-phase precipitate is formed at the expense of dissolving approx. 1000–1500 finely dispersed MX precipitates in the matrix. The disappearance of MX precipitates in martensitic steels during service in favor of Z-phase eliminates the effect of precipitation hardening with these particles. Nevertheless, as shown in [12], the interaction of MX precipitates with dislocations (Figure 9b) is still observed in the microstructure of P91 steel after service, and single Z-phase precipitates do not adversely affect its properties, and thus the creep strength (Figure 10).
Figure 10.
Results of short-term creep tests of T91 steel after service [12].
Unlike the MX precipitates, both the chemical composition and the size of Z-phase depend on chemical composition of the steel it precipitates in and on creep duration. The Z-phase in 9%Cr steels is approx. 80–100 nm, whereas in steels with 11–12%Cr it is much larger and amounts to approx. 0.5–2 μm. Consequently, in 9%Cr steels the Z-phase precipitation is accompanied by a slight reduction in the volume fraction of MX precipitates, whereas in 12%Cr steels MX precipitates are virtually completely transformed into this complex nitride [31, 32, 36]. In addition, in 9%Cr steels the Z-phase precipitates after approx. 105 h at the earliest, while in 12%Cr steels the precipitation of this phase can be observed as early as after 103 h. Hence, the effect of Z-phase precipitates on creep strength is slight in steels with 9%Cr, whereas in 12%Cr steels it is significant [25, 32, 35, 36].
In high-chromium martensitic steels, the Z-phase precipitation may proceed with two mechanisms [32, 33]. The schematic transition of MX precipitates into Z-phase in high-chromium martensitic steels is shown in Figure 11.
Figure 11.
Schematic transition of MX precipitates into Z-phase [32].
On the other hand, dissolving NbX precipitates in the matrix “releases” carbon atoms, which results in the precipitation of chromium-rich M23C6 carbides, frequently nearby the Z-phase particles. The precipitation of Z phase is preferential near the grain boundaries of prior austenite, and in the steels containing delta ferrite additionally also near the interphase boundary martensite/delta ferrite [14, 25, 33, 35, 36]. This is due to faster diffusion of substitution elements nearby these defects. It results in the formation of near-boundary areas free from MX precipitates, which leads to the accelerated matrix recrystallization and reduction in strength properties in these areas. Such changes lead to the formation of creep grain, which is unequal in volume, and consequently to a faster destruction of steel during service [32, 35, 36]. The disappearance of finely dispersed MX precipitates also results in disproportionately high reduction in hardness in relation to other mechanical properties [21].
In martensitic steels containing approx. 9%Cr, a more important problem is the increase in the size of M23C6 carbides as well as the precipitation and growth of intermetallic Fe2Mo Laves phase [6, 7, 25, 37]. In 9–12%Cr steels in the as-received condition, the Laves phase does not occur. The precipitation of this phase takes place during service/aging mainly at the grain/lath boundaries, frequently nearby M23C6 carbides Figure 12 [37–39]. In the case when the total content of W + Mo in the steel amounts to at least 4.53, the particles of Laves phase precipitate heterogeneously at grain boundaries as well as homogeneously within grains, forming the precipitation free zone on both sides of the grain boundary [40].
Figure 12.
Precipitation of Laves phase at the grain boundary nearby M23C6 carbides: (a) T91 steel, (b) GP91 cast steel [38].
It is assumed that due to high dispersion in the initial period of the precipitation the Laves phase has a positive effect on properties of these steels by increasing the precipitation hardening. However, low stability of the Laves phase results in its high coagulability, which results in a very fast increase in its size [15, 25, 37, 39].
The Laves phase precipitating in 9–12%Cr creep-resistant steels makes the matrix deplete of substitution elements (tungsten, molybdenum, chromium), which increases the tendency of these steels to the recovery and polygonization process and reduces their resistance to oxidation. On the other hand, the matrix depletion of substitution elements (Cr, Mo, W), which are also components of M23C6 carbides, has a positive effect on the inhibition of coagulation of these precipitates [12, 38]. According to [41, 42], the nucleation and growth of Laves phase requires the enrichment of micrograin boundaries not only in Mo and Si, but also in phosphorus.
The Laves phase precipitates with average diameter above 130 nm also contribute to the change in cracking mechanism from ductile to brittle (transcrystalline, cleavable fracture) and are the main reason for sudden reduction in creep strength of [40, 41, 43, 44]. The Laves phase and M23C6 carbide precipitates occurring during long-term service form the so-called continuous grid of precipitates at the grain boundaries (Figure 13), which contributes to a decrease in ductility of 9% Cr steel [8, 12, 43–49].
Figure 13.
Continuous grid of precipitates at the former austenite grain boundaries (a) P91 steel (SEM), (b) GP91 cast steel (TEM) [49].
The increase in stability of the Laves phase precipitates in these steels can be achieved by the addition of boron and/or tungsten [50, 51]. On the contrary, phosphorus, silicon, and cobalt have an adverse effect as they stimulate the precipitation of the Laves phase [52, 53].
In the tempered structure of martensite in steel of the 9–12%Cr type, in the case of the typical volume fractions of particular precipitates and spacing between them, the Orowan stress can be estimated as shown in Table 2.
Particle
Volume fraction, %
Diameter, nm
Spacing, nm
Orowan stress, MPa
Fe2M
1.5
70
410
95
M23C6
2
50
260
150
MX
0.2
20
320
120
Table 2.
Volume fraction, diameter, and spacing of each kind of precipitates in high-chromium martensitic steel, together with Orowan stress from the values of interparticle spacing [14].
The precipitation processes as well as the growth of carbides and secondary phases, which occurs during the operation of creep-resistant steels, make the matrix deplete of substitution elements as a result of their diffusion into these precipitates.
The matrix depletion of the above-mentioned elements facilitates the self-diffusion processes, speeds up the matrix recovery and polygonization processes, and reduces the oxidation resistance, thus contributing to the reduction in high-temperature creep resistance and life of these steels [11, 54].
An important factor affecting the basic property of 9–12%Cr steels, i.e., creep strength, is aluminum content in the steel. The addition of aluminum to steel is to deoxide it in the metallurgical process, hence part of aluminum will remain as Al2O3 in the steel, while the other part in the atomic form is dissolved in solid solution. The aluminum content in 9–12%Cr steels for the power industry should not exceed 0.04%. Due to its greater affinity to nitrogen, aluminum forms large lamellar AlN nitrides of about 0.5–1.0 μm to prevent the formation of finely dispersed vanadium-rich MX precipitates. This results in not only a reduction in the volume fraction of vanadium-rich MX precipitates, but also in a change in their chemical composition. There occurs a decrease in the amounts of vanadium and nitrogen with simultaneous increase in the amounts of niobium and carbon in these precipitates, which makes niobium-rich carbides (NbC) become the main precipitate in martensitic steel. This results in a decrease in creep strength of martensitic steels: for P91—by about 10% at 600°C, whereas for P92 the time to rupture of test specimen was shorter by 7–30 times, depending on test temperature and stress [30, 55]. Higher than permissible aluminum content in 9–12%Cr steels has a positive effect on the increase in impact strength and reduction in brittle fracture appearance transition temperature, which is the result of positive effect of AlN on the austenite grain refinement. The increase in aluminum content from 0.03 to 0.094% in P92 steel makes the average diameter of austenite grain decrease from 50 μm (which corresponds to the grain grade of 5.5) to 10 μm (grain grade 10) [30, 56]. The negative impact of aluminum on creep resistance requires control of chemical composition of 9–12%Cr steel as early as at its production stage, in particular with regard to elements with high affinity to nitrogen, such as Al and Ti, so as to prevent from the formation of unfavorable AlN or TiN nitrides at the expense of finely dispersed VN precipitates.
The basic features of the microstructure of 9–12%Cr steel in the as-received condition and after long-term service/aging are summarized in Table 3.
Features of microstructure
As-received condition
Creep/aging
Matrix
Dislocation density
High
Low/very low
Size, width of subgrains/martensite laths
High-temperature tempered martensite microstructure with small width of laths
Recovery and polygonization process – transformation of the lath martensite microstructure into the polygonised ferrite grain microstructure
Precipitates
MX
Finely dispersed (~20–50 nm), precipitated inside laths at dislocations, limit grain growth, make steel precipitation-hardened
Finely dispersed (~20–50 nm), precipitated inside laths at dislocations, limit grain growth, make steel precipitation-hardened, change into Z-phase
M23C6
50–150 nm, precipitated at the martensite lath boundaries and at the former austenite grain boundaries, stabilize substructure
≥220 nm, partially precipitated at the subgrain boundaries, coagulate during creep/aging resulting in a reduction in creep strength and increase in embrittlement
Z-phase
Nonexistent
Formed at the expense of finely dispersed MX precipitates, causes sudden decrease in creep strength
Laves phase
nonexistent
Medium and large size of precipitates (≥0.5–1 μm), precipitated at the grain and subgrain boundaries nearby M23C6, decreases creep strength and reduces ductility
Table 3.
Basic features of microstructure in 9–12%Cr steels in the as-received condition and after creep.
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Written By
Grzegorz Golański, Cezary Kolan and Joanna Jasak
Submitted: 10 April 2017Reviewed: 08 August 2017Published: 20 December 2017
Open access peer-reviewed chapter
