Creep properties of simulated heat-affected zone of HR3C austenitic steel

Materials Characterization 128 (2017) 238–247

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Creep properties of simulated heat-affected zone of HR3C austenitic steel a,b,⁎

V. Sklenička a b

a

, K. Kuchařová , M. Kvapilová

a,b

, M. Svoboda

a,b

, P. Král

a,b

, J. Dvořák

a,b

MARK

Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Žižkova 22, 616 62 Brno, Czech Republic CEITEC – Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Žižkova 22, 616 62 Brno, Czech Republic

A R T I C L E I N F O

A B S T R A C T

Keywords: HR3C steel Welding Heat affected zone Creep Microstructure Fractography

Tensile creep tests on austenitic stainless HR3C steel pipe in an as-received state and after heat affected zone (HAZ) thermal cycle simulation by means of the GLEEBLE 3800 physical simulator were carried out at 923 and 1023 K and an applied stress range between 100 and 350 MPa to evaluate the effect of welding on its creep properties and behaviour. It was found that creep testing was conducted in a power-law creep regime. The results show a clear detrimental influence of the HAZ simulation on creep properties of HR3C steel at 923 K and for very short creep exposure. However, with increasing time to fracture, the effect of weld simulation becomes less significant. By contrast, no such deterioration of creep properties was detected at 1023 K. The evaluated similar values of the stress exponents of the minimum creep rate and the time to fracture are explained using the assumption that both creep deformation and the fracture process are controlled by the same mechanism. The main creep fracture mode was intergranular brittle fracture due to grain boundary creep damage.

1. Introduction Advanced austenitic stainless steel grades are highly valued for fabrication of high temperature components in ultra-supercritical (USC) thermal and nuclear power generation plants [1–4]. Among these, HR3C austenitic steel has become a widely used material for superheater and reheater tubes in USC units [5]. HR3C steel is an improved version developed from a 25Cr-20Ni series steel of SUS310S. Thus, HR3C steel possesses not only excellent resistance to high-temperature corrosion and steam oxidation due to high Cr content but also superior creep strength [6–9]. Nitrogen, as an alloying element, acts as a creep strengthening agent in 25%Cr austenitic matrix and the combination of 0.4%Nb and 20%Ni additions is an effective measure of the enhancement of creep strength. Most previous studies on HR3C steel focused mainly on the impact of the alloying elements and some individual precipitated secondary phases. Relatively limited studies are conducted on creep behaviour of HR3C steel and in particular its welded joints [10–12]. Welding is still the major joining and repair technology for power plant components [13,14]. Long-term experience of welded structures exposed to creep shows that the heat affected zone (HAZ), a narrow zone of base material adjacent to the weld fusion line altered by the weld thermal cycle, in respect to the creep strength, is often regarded as the weakest link in welded construction [15]. In the last two decades, an intensive effort was devoted to understanding the creep behaviour and properties of weldments of advanced 9–12%Cr ferritic steels ⁎

Corresponding author at: IPM AS CR, Zizkova 22, 616 62 Brno, Czech Republic. E-mail address: [email protected] (V. Sklenička).

http://dx.doi.org/10.1016/j.matchar.2017.04.012 Received 11 November 2016; Received in revised form 3 April 2017; Accepted 11 April 2017 Available online 12 April 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.

[15–22]. By contrast, a limited effort was given to creep-exposed austenitic weldments. One of the major problems for austenitic heatresistant steels is the reduced creep strength of the welded joint during low-stress and long-term creep deformation. The welded joint is normally fractured in the HAZ [23–29].Thus, research into the evaluation of the HAZ microstructure during creep and its implications for creep damage behaviour in heat-resistant austenitic steels is vital in order to attain a fundamental understanding of the mechanisms of creep deformation and fracture. The problem of conducting microstructural and creep investigations of the HAZ in real welds is the presence of extremely small and inhomogeneous HAZ. However, HAZ simulation allows the study of uniform microstructure and properties that represent HAZ. This generated homogeneous microstructure can be used for all kinds of microstructural investigations and creep testing. Different techniques are currently applied for HAZ simulation; Buchmayer [30] provides a good overview. One of the frequent types of HAZ simulation is controlled resistance heating of the specimen using a weld simulator GLEEBLE™ [15,31,32]. This study aims to investigate the creep properties and the corresponding microstructural changes of the simulated HAZ of HR3C steel. HAZ thermal cycle simulation was performed using the GLEEBLE 380 physical simulator. The creep tests in tension were performed at two testing temperatures. To complete the characterisation of creep behaviour and a microstructural evolution in the HAZ, the concurrent creep tests were run on the HR3C steel in the as-received (unsimulated)

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state. The creep tests were followed by microstructural characterisation and fractographic analyses of the tested specimens by means of light metallography and scanning and transmission electron microscopy.

Table 1 Results of tensile tests of AR and simulated HAZ states of HR3C steel. Temperature

Material

Mechanical properties

T [K]

Location

Rp0.2 [MPa]

Rm [MPa]

A10 [%]

Z [%]

293 293 873 873 873 873

AR HAZ AR AR HAZ HAZ

350 348 189 194 248 229

755 751 495 502 522 506

49.0 70.0 48.3 43.0 62.0 62.5

73.3 39.9 75.0 75.0 72.4 75.0

2. Material and Experimental Procedures 2.1. HR3C Steel, Processing and Tensile Testing The HR3C steel was received in the as-received state after solution treatment in the form of the pipe. The dimensions of this pipe were Outside Diameter 45 mm × Wall Thickness 10 mm. Its chemical composition was as follows (in wt%): 0.061 C, 25.47 Cr, 20.73 Ni, 0019 P, 0.007 S, 0.35 Nb, 0.146 Mo, 1.109 Mn, 0.52 Si, 0.20 N and Fe balance. With reference to the ASTM standard requirement, the chemical composition of the investigated HR3C steel is within the maximum and minimum range. The pipe was solution annealed at 1503 K for 30 min in order to adjust the grain size and to dissolve the precipitates which had formed during manufacturing of the pipe. The as-received state (without simulation) of the HR3C steel pipe will be denoted as the (AR) condition. A weld HAZ thermal cycle Tmax = 1573 K and Δ t8/ 5 = 60 was simulated on samples of HR3C steel by means of GLEEBLE 3800 physical simulator, which enlarged the characteristic regions of the HAZ to study the microstructure and creep properties. The GLEEBLE simulation was performed at Welding Research Institute – Industrial Institute of the Slovak Republic, Bratislava, Slovakia. A typical record of the simulated HAZ thermal cycle is illustrated in Fig. 1. After the thermal cycle a post-welding heat treatment (PWHT) was carried out at 1073 K for 4 h, followed by a cooling in a switched off furnace followed by further cooling in air. The state of the steel after the HAZ thermal cycle simulation will be denoted as the (HAZ) condition. The dimension of the specimens used in the GLEEBLE simulator was Ø 6 mm × 85 mm and the simulated characteristic zone was set as a central reduction of Ø 3 mm × 15 mm in the middle of the specimen. Tensile properties of the pipe for both material states at room temperature and at 873 K temperature are presented in Table 1. No strengthening of the HAZ results is observed in Table 1. In line with tensile properties the hardness measurement (HV10) of the AR and the HAZ at room temperature did not show any significant difference. The mean values of HV10 for states AR and HAZ were 173 and 174, respectively.

2.2. Creep Testing Constant load creep tests in tension were carried out in an argon atmosphere until the final fracture of the specimens. Creep specimens were machined from the middle parts of the wall thickness of the pipe. The creep testing was conducted at 923 and 1023 K with the testing temperature maintained to within ± 0.5 K of the desired value. The initial applied stresses ranged from 100 to 350 MPa. The creep elongations were measured using a linear variable differential transducer (the strain was measured with a sensitivity of 5 × 10− 6) and they were continuously recorded digitally and then computer processed [22]. 2.3. Microstructural and Fractographic Investigations The crept specimens were prepared for microstructural and fractographic examinations by means of light metallography and scanning and transmission electron microscopy. Transmission electron microscopy (TEM) studies were carried out on carbon replicas and thin foils prepared from both simulated and nonsimulated parts of the crept specimens using a Jeol 2100F electron microscope operating at 200 kV and equipped with X-Max80 Oxford Instruments EDS detector for X-ray microanalysis. Particles of secondary phases extracted into carbon replica were identified by means of selected area diffraction (SAED) and their local chemical composition was measured by energy dispersive X-ray spectroscopy (EDS). The creep fracture surfaces and the creep damage and fracture profiles on longitudinal metallographic sections of fractured creep specimens were investigated using a Tescan Lyra 3 scanning electron microscope (SEM). 3. Experimental Results 3.1. Creep Behaviour The results of creep tensile tests carried out at two different testing temperatures on the AR and simulated HAZ state steel HR3C are summarised in Tables 2 and 3, respectively. Comparing the data relevant to creep lives of different the material states, substantial differences in their creep behaviour can be found at 923 K. By contrast, no such essential differences can be generally noticed at 1023 K. Figs. 2 and 3 display the standard creep curves of HR3C steel in the Table 2 Results of creep testing on AR state of HR3C steel.

Fig. 1. Simulated weld HAZ thermal cycle Tmax = 1573 K, Δt 8/5 = 60 s.

239

Temperature T [K]

Stress σ [MPa]

Time to fracture tf [h]

Minimum creep rate ε̇m [s− 1]

Elongation εf [%]

923 923 923 923 1023 1023 1023

220 250 300 350 100 150 200

2079.841 1016.675 231.061 23.685 1829,842 212.449 32.963

2.2 × 10− 9 1.7 × 10− 9 1.1 × 10− 8 8.6 × 10− 8 4.2 × 10− 9 7.6 × 10− 8 2.7 × 10− 7

13.2 11.0 10.1 14.0 10.1 30.1 31.7

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0.20

Table 3 Results of creep testing on simulated HAZ state of HR3C steel.

873 923 923 923 923 1023 1023 1023 1023 1023 1023

Stress σ [MPa]

Time to fracture tf [h]

300 220 250 300 350 100 120 150 180 200 250

367.474 1189.762 361.13 52.135 1.557 1713.420 634.414 168.533 56.026 34.870 6.748

Minimum creep rate ε̇m [s− 1] 5.5 × 10− 9 1.5 × 10− 9 3.6 × 10− 9 4.6 × 10− 8 1.3 × 10− 7 1.2 × 10− 8 2.6 × 10− 8 5.9 × 10− 8 1.3 × 10− 7 1.6 × 10− 7 1.0 × 10− 6

steel HR3C-HAZ

Elongation εf [%]

923 K 0.15

13.7 6.5 8.4 13.2 17.6 18.3 22.9 18.9 16.7 15.8 14.1

STRAIN ε

Temperature T [K]

300 MPa 0.10

250 MPa σ = 220 MPa

0.05

(a)

0.20 0.00

steel HR3C σ = 220 MPa 923 K

0.10

0.00

400

600

800

1000

1200

0.25

150 MPa

0.20

HAZ

0.05

200

TIME t [h]

STRAIN ε

STRAIN ε

0.15

0

AR

ocel HR3C, HAZ 1023 K

120 MPa

0.15

0.10 σ = 100 MPa

0

500

1000

1500

2000

2500

0.05

TIME t [h]

(b)

Fig. 2. Standard creep curves of strain vs. time for HR3C steel in AR and HAZ states at 923 K and 220 MPa.

0.00

0

500

1000

1500

2000

TIME t [h] 0.20

Fig. 4. Comparison of the values of the instantaneous strain ε0 for the HAZ state and different applied stress σ at (a) 923 K and (b) 1023 K.

steel HR3C

STRAIN ε

0.15

0.10

HAZ

0.05

0.00

AR and the HAZ states for the tests conducted under constant tensile loads at temperatures of 923 and 1023 K. These standard creep curves graphically represent the time dependence of strain measured over the whole gauge length (the AR condition) or within the HAZ zone only (simulated HAZ condition) of creep specimen during the creep exposure. As demonstrated by Figs. 2 and 3, significant differences were found for HAZ state material creep behaviour at two testing temperature conditions. First, all standard ε vs. t curves appear to show the presence of instantaneous strain ε0 upon loading, which contains elastic, anelastic and plastic strain components [33,34]. The values of ε0 are considerably higher at a temperature 923 K (Fig. 4a) than those at 1023 K (Fig. 4b) and at 923 K strongly depend on the stress level (Fig. 4a). By contrast, the values ε0 for a temperature of 1023 K seem to be practically independent of the applied stress (Fig. 4b). Second, the shapes of the creep curves for temperatures 923 K and 1023 K differ considerably. Third, the values of the strain to fracture εf are generally higher at a temperature of 1023 K and perhaps more, depending on stress at 923 K. It is important to note that the creep curves shown in Fig. 4 are dominated by the tertiary creep stage. Further differences in the creep behaviour of the HR3C steel in AR and HAZ states loaded under the same conditions are confirmed and

σ = 100 MPa 1023 K

AR

0

500

1000

1500

2000

TIME t [h] Fig. 3. Standard creep curves of strain vs. time for HR3C steel in AR and HAZ states at 1023 K and 100 MPa.

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-3

10

. CREEP RATE ε [s-1]

. MINIMUM CREEP RATE ε m [s-1]

steel HR3C

-4

10

923 K σ = 220 MPa

-5

10

-6

10 10

HAZ

-7

AR

-8

10 10

-9

-10

10

10

-5

10

-6

10

-7

10

-8

10

-9

Steel HR3C

AR HAZ

923 K 1023 K

(a) -10

0

500

1000

1500

2000

10

2500

100

1000

Fig. 5. Replotted standard creep curves for the AR and HAZ states shown in Fig. 2.

-2

Steel HR3C

10

1023 K σ = 200 MPa

-4

STRESS σ [MPa]

steel HR3C

-3

10

. CREEP RATE ε [s-1]

1000

STRESS σ [MPa]

TIME t [h]

AR

10

HAZ -5

10

AR HAZ

500

923 K 1023 K

300

-6

10

(b) 100

-7

10

1

100

1000

10000

TIME TO FRACTURE t [h]

-8

10

10

0

10

20

30

40

f

50

Fig. 7. Stress dependences of (a) the minimum creep rate ε̇m and (b) the time to fracture tf.

TIME t [h] Fig. 6. Replotted standard creep curves for the AR and HAZ states at 1023 K and 200 MPa.

in stress on the creep rate appears in the tertiary creep stage. The creep data of both states of the HR3C steel are summarised in Fig. 7, where the minimum creep rate ε̇m and the time to fracture tf are plotted against the applied stress σ on a double logarithmic scale. Inspection of Fig. 7a leads to two observations. First, the values of the minimum creep rate at 923 K are mostly lower for state AR than those for HAZ. By contrast, no clear tendency was found for temperature 1023 K. Second, the slopes and therefore the values of the apparent stress exponent of the creep rate n = (∂ ln ε̇m / ∂ ln σ)T are similar for both states. The double logarithmic plots of the time to fracture tf as a function of applied stress are shown in Fig. 7b. It is clear from these plots that the creep life of the AR state at 923 K can be an order of magnitude longer at high stresses than that for the simulated HAZ state. However, this difference consistently decreases with decreasing applied stress. On the other hand, no substantial differences in creep lives were found between AR and HAZ states at 1023 K. The evaluated values of stress exponent m of the time to fracture tf (m = ∂ ln tf / ∂ ln σ)T derived from Fig. 7b do not differ significantly for both the states of material.

perhaps more clearly illustrated by relevant standard creep curves displayed in Figs. 2 and 3. Generally, the AR state exhibits better creep resistance than the steel in HAZ state at a temperature 923 K. Whereas the values of the instantaneous strain ε0 are higher and the time to fracture tf are shorter for HAZ state at a temperature 923 K; although no substantial differences in those parameters between AR and HAZ were found at 1023 K. It should be stressed that the standard creep curves shown in Figs. 2 and 3 do not clearly indicate the individual stages of creep. However, these standard ε vs. t creep curves can be easily replotted in the form of the creep rate ε̇ vs. time t as shown in Figs. 5 and 6. It is clear that neither curve exhibits a well-defined steady state. In fact, this stage is reduced to an inflection point of the ε̇ vs. t curve. The creep rate of both states decreases significantly during the primary creep stage until reaches minimum, which is then followed by creep acceleration. The minimum creep rate ε̇m is reach very early after loading. The curves show that the primary stage is very short and an extensive tertiary stage covers practically the whole duration of creep exposure. It should be note that under the constant tensile load, the stress continuously increases as creep damage proceeds or as the crosssection of a specimen decreases and a pronounced effect of the increase

3.2. Microstructural Investigations According to already published results of detailed microstructural 241

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Fig. 9. TEM micrograph of coarse Z-phase particle (AR state, crept at 923 K, 350 MPa) and corresponding SAED patterns of the zone axis [100].

only three types of precipitates were unambiguously identified as M23C6 of the type (Cr,Fe)23C6, coarse Z-phase (NbCrN) and fine Nb (N,C) precipitates in both structural states of the steel. σ-Phase and Cr3Ni2Si(C,N) precipitates were not found in any state, evidently due to short creep exposures [11,37]. M23C6 carbides, having elongated spherical morphology, were mainly located at grain boundaries (Fig. 8a) but also close to grain boundaries and within grains as cuboidal precipitates (Fig. 8b) and/or as long plates parallel to incoherent twin boundaries (Fig. 8c). Coarse round particles of Z-phase were distributed in the grain interior (Fig. 9). Very fine Nb(C,N) precipitates were found to nucleate intragranularly as well as along grain boundaries (Fig. 10). However, the density of Nb(C,N) precipitates in the AR state was markedly lower than in the HAZ state before creep exposures. It should be noted that Sourmal and Bhadeshia [39] consider these precipitates as residual particles that formed during solidification and are undissolved by the solution treatment. By contrast, Tassa et al. [8] reported that Z phase was often formed by nucleating and growing on niobium carbonitrides. They explained this observation by considering that Nb(C,N) precipitates, formed at high temperatures during fabrication process, are unstable at the exposition/ ageing temperatures and they are transformed into stable Z phase nitride which less contribute to the creep strength [8]. This could explain why the difference in the creep strength between the AR and HAZ states decreases with an increase in time of creep exposure at 923 K and was not observed at 1023 K. The large volume fraction of these precipitates suggests the generation of dislocations by the growing particles, which have a smaller atomic volume than the matrix (Fig. 10). Fig. 8. TEM micrographs of M23C6 carbides in the AR state located at (a) grain boundaries (crept at 923 K, 350 MPa), (b) within grains (crept at 923 K, 220 MPa) and (c) at twins (crept at 923 K, 220 MPa). Inlays show SAED patterns from austenite (strong spots) and M23C6 (weak spots, indexed in italics) of the zone axes (a) [110] and (b) [100].

SEM and TEM investigations of aged and/or crept HR3C steel specimens the presence of six different precipitates was revealed [7–12,35,36]: M23C6, Cr2N, σ-phase, Z-phase, η-phase (Cr3Ni2Si(C,N) and Nb(C,N)). These precipitates were predicted and confirmed by MatCalc simulation [8,9,37,38]. In this work microstructural investigation as a supplement of these previous investigations focused mainly on such microstructural features that could account for differences between the AR state and the simulated HAZ state. Light microscopy analysis revealed that differences in grain size of the AR and simulated HAZ states before creep exposure is relatively small. Coarse grain sizes for the AR and the HAZ states were 99 μm and 135 μm, respectively. This difference is associated with the different thermal history of the HAZ state. However, during creep exposures further small coarsening of grain size occurred. Using TEM and SEM

Fig. 10. TEM micrograph showing dislocations punched out at fine Nb(C,N) precipitates in the grain interior (AR state, 1023 K, 200 MPa).

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Fig. 11. A widening of grain boundaries during creep in the AR state: (a) creep exposure 23.7 h (923 K, 350 MPa), (b) creep exposure 2079.9 h (923 K, 220 MPa). Fig. 13. Composition profile along the corresponding line shown in the TEM image (HAZ, creep at 1023 K, 220 MPa).

In contrast to previously published reports, a yet undocumented microstructural change that could affect the creep behaviour of HR3C steel is revealed in this study. Microstructural investigation by SEM after creep testing of both the AR and the HAZ states showed a considerable increase in the thickness of grain boundaries during creep exposure. This effect is clearly demonstrated by Figs. 11 and 12, which show considerable widening of grain boundaries when compared to similar observations at shorter creep exposures. In the case of very short-term creep testing practically no such widening was observed (Figs. 11a and 12a). Conversely, after longer creep exposures the “multilayer” grain boundaries exist consisting of M23C6 particles and the narrow regions near grain boundaries are substantially free of precipitates and/or contain precipitates in reduced amounts of very fine particles of Nb (C,N) (Figs. 11b and 12b). Chemical composition measured by EDS line analysis across such a grain boundary is shown in Fig. 13. The “multilayer” boundaries are widest in the HAZ and less extensive in the AR state. Analysis of the TEM micrographs shows that the dislocation substructure of the AR state was qualitatively similar to the HAZ one, however, dislocation density seems to be higher in the case of the HAZ state than the dislocation density in the matrix of the AR state (Fig. 14). Interactions of free dislocation, dislocation pile-ups and the work hardening zone (WHZ) characterised by a dislocation density near the grain boundary were frequently observed (Fig. 15). By analyzing the dislocation substructure the nucleation sites of Nb(C,N) particles were considered to be dislocations (Fig. 10). 3.3. Fractographic Investigations

Fig. 12. A widening of grain boundaries during creep in the HAZ state: (a) creep exposure 1.56 h (923 K, 360 MPa), (b) creep exposure 1189.8 h (923 K, 220 MPa).

After the creep tests, the fractured creep specimens were analyzed. 243

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Fig. 14. TEM micrograph of the dislocation network in the AR state (creep at 1023 K, 200 MPa).

Fig. 16. SEM micrographs of creep fracture surfaces of the AR and HAZ states: (a) AR state, creep at 923 K and 220 MPa, (b) HAZ state, creep at 923 K and 220 MPa.

Fig. 15. TEM micrograph showing the interaction of dislocations (AR, creep at 1023 K, 200 MPa).

Fracture surface examination by SEM of crept specimens at various applied stresses, two testing temperatures and both states of the steel showed that the main mode of creep fracture was intergranular with cracks and/or cavities along grain boundaries. Distinctive flat grain boundary facets (Fig. 16) covered the dominant areas of the creep fracture surfaces in all of the tested specimens, demonstrating intergranular mode of creep fracture. These findings were supported by observations of creep damage and fracture on metallographic longitudinal sections of fractured specimens (Figs. 17 and 18). Typical intergranular creep fracture occurs by loss of internal cross section due to intensive intergranular cavitation [16,40–42]. Creep cavities and/or microcracks were observed predominantly in the vicinity of the fracture (Figs. 17 and 18). In the higher stress regime cavities are formed frequently at triple grain junctions, like wedge cracks [41]. In the lower stress regime cavities have a rounder shape. However, independent of their shape or morphology, cavities are usually attached to grain boundary coarse particles of M23C6.

4. Discussion 4.1. Creep Behaviour Upon Loading As demonstrated by Figs. 2–4 significant differences were found in the creep behaviour of the steel at temperatures of 923 and 1023 K. The instantaneous strain ε0 upon loading of creep tests at a temperature 923 K represents a prevailing contribution to the fracture strain εf, especially at higher stresses and for the HAZ state; whereas at 1023 K the strain ε0 practically disappeared and the fracture strain εf is

Fig. 17. Creep fracture profiles on metallographic longitudinal sections: (a) AR state, creep at 923 K and 220 MPa, (b) HAZ state, creep at 923 K and 220 MPa.

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The temperature dependence of the minimum creep rate reflects the fact that creep involves thermally activated processes. To identify the process that controls the creep rate, knowledge of the activation energy for creep Qc is required [33]. The activation energy Qc can be estimated as Qc = [∂ ln ε̇m / ∂ (−1/kT)]σ, where k is the Boltzmann's constant. Using Tables 2 and 3, determined values of Qc range from 293 to 430 kJ/mol, increasing with stress for the HAZ state. It should be noted that Park et al. [6] determined values of Qc ≈ 330 kJ/mol (theoretical approach) and Qc ≈ 315 kJ/mol (experiment) for HR3C steel crept in the temperature range of 923 to 973 K. These values of Qc are consistent with those of the lattice self-diffusion. According to Vujic et al. [9], solid solution raises the activation energy for creep Qc. This may explain why the determined values of activation energy Qc of austenitic stainless steel are considerably larger than that for selfdiffusion. Generally, three main contributions to the creep strength of HR3C steel should be considered: solid solution hardening, precipitation hardening and dislocation hardening. Very recently, Vujic et al. [9,37] used thermo-kinetic calculations of the precipitation structure in the austenitic stainless steel 25Cr20NiNbN with the help of the software MatCalc [45,46]. The results were used for creep strength predictions at 923 and 1023 K [9]. It was found that the dislocation hardening, followed by precipitation hardening, gives the largest contribution. In spite of the fact that Nb has a large lattice misfit parameter, the solid solution hardening contribution is obviously small for the investigated steel.

Fig. 18. Intergranular creep damage in the AR state (creep at 923 K and 220 MPa).

dominated by creep strain originated during the tertiary creep stage. It can be assumed that the applied stresses σ at 923 K are higher than the thermal yield stress σa [34]. Thus, in the case that σ ≥ σa, dislocations are introduced during fast plastic deformation upon loading and the dislocation structure at the beginning of creep tests is different from that when the applied stresses σ is below σa. Different values of ε0 during tests of AR and HAZ states, which were carried out under the same loading conditions, may be caused by different arrangements of dislocation structures in both states before the test as a result of their different heat treatment history. 4.2. Creep Mechanisms

4.3. Creep Damage and Fracture

In analyzing creep behaviour of the steel under investigation, two questions naturally arise. The first question is in which creep regime the creep tests were performed? The second question is which creep mechanisms operate in this regime? The answers to the above questions should be provided by activation analysis of creep data. As shown in Figs. 5 and 6, immediately after the primary creep the tertiary creep stage begins. In this case, the minimum creep rate ε̇m can be defined instead of the steady-state creep rate ε̇s. The minimum creep rate ε̇m can be explained by the process whereby hardening in the primary stage is balanced by softening in the tertiary stage [33]. The values of the stress exponent of the minimum creep rate n = (ln ε̇m / ln σ)T can be estimated as the slopes of the plots presented in Fig. 7a. For the AR state the values of n of 8.3 and 6.1 were found for at temperatures of 923 K and 1023 K, respectively, whereas for the HAZ state the value n of 10.2 and 5.1were found for a temperature of 923 K and 1023 K, respectively. These values of stress exponent n are high compared to those predicted by dislocation climb controlled models of power-law creep [33] but do not comply to the criterium proposed by Sherby and Burke [43] for the power-law stress region break down. However, these values are in the range typical for precipitationstrengthened materials in a high stress regime, whereby tertiary creep governs creep life [6] and are comparable to reported values in the literature. Further, referring to Lagneborg's creep model of precipitation-strengthened alloys [44] a gradual increase of the stress exponent n may be due to an overlapping of the stress intervals characteristic for the dislocation climb and the Orowan mechanism. The high values of the stress exponent n probably involve internal stresses in addition to a power-law dependence of the flow rate with respect to an effective stress. Thus, the stress dependence of the minimum creep rate obeyed Norton's power law whereby creep behaviour is controlled by the climb of dislocations [33]. The evaluated values of the stress exponent m of the time to fracture (m = −(∂ ln tf / ∂ ln σ)T) (as the slopes in Fig. 7b) are similar to the values of the minimum creep rate stress exponent n (see Section 3.1). This can be explained by the assumption that both the creep deformation and fracture are controlled by the same mechanism (s).

As mentioned in Section 3.3, fractographic investigations indicate that the main fracture mode was intergranular brittle creep fracture due to cavities and/or microcracks situated along grain boundaries. Preferred nucleation sites for cavities were particle/matrix interfaces associated with grain boundary second phase particles like M23C6. Very recently, Sandström and coworkers [42,47,48] analyzed and proposed the concept of creep cavitation in austenitic stainless steels, based on the postulate that the cavity nucleation is assumed to take place where subboundaries on one site of a sliding grain boundary meet subgrain corners on the other site (the double edge model [49]). Alternative cavitation positions can be found where particles meet subboundary [42]. It should be emphasised that grain boundaries are densely decorated by grain boundary particles (Fig. 8), which could effectively inhibit grain boundary sliding. As shown in Fig. 14, the work hardening zone consists of accumulated dislocations and their stress fields are the direct microstructural source of backstresses created in the matrix. The structure of the work hardening zones would not be stable against some recovery mechanism(s). Such accommodation and recovery mechanisms may be damage processes, like nucleation cavities by means of debonding (decohesion) of the grain boundary particle-matrix interface. Consequently, more intensive recovery decreases efficiently with dislocation density in the work hardening zone and as a consequence, the stress transferred to the particle, as well as the back stress imposed into the matrix from the work hardening zone, decreases. Further cavity growth and coalesce cavities are obviously controlled by plastic flow, which is consistent with the similarity of the values of the stress exponents n and m (Section 4.2). The final stage of intergranular creep fracture occurs due to the critical decrease of the grain boundary cohesion caused by the critical accumulation of cavitation damage in the whole specimen volume. The critical cavitation damage can be defined as a state of damage, at which the threshold probability of damage interaction over distances comparable with the specimen cross-section dimension is attained [40,41].

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1000 AR

Steel HR3C

creep rate n and the time to fracture m were explained using the assumption that both creep deformation and fracture processes are controlled by the same mechanism(s). The main creep fracture mode was intergranular brittle fracture due to the formation and accumulation of grain boundary creep damage. 4. The predicted creep lives were calculated using the ECCC master equation for the steel under investigation [43]. The comparison of the experimental points with those calculated according the master equation revealed that in the AR state the predicted and experimentally determined times to fracture followed an excellent correlation at both testing temperatures. However, although the simulated HAZ specimens also closely followed the same one-to-one correlation at 1023 K, they failed at 973 K significantly earlier than was the case for the AR material and the prediction.

HAZ

STRESS σ [MPa]

973 K 1023 K ECCC (AR)

973 K 1023 K

Acknowledgements 100 1

10

100

1000

TIME TO FRACTURE tf [h]

Part of this work was sponsored by the European Union (directorate-general for energy), within the project MACPLUS (ENER/FP7EN/ 249809/MACPLUS) in the framework of the Clean Coal Technologies. The authors would like to recognize Dr. Peter Bernasovský of Welding Research Institute – Industrial Institute of the Slovak Republic, Bratislava, for his support of HAZ GLEEBLE simulation.

10000

Fig. 19. The comparison of predicted creep lives using ECCC master equation with experimental values.

References

4.4. Predicted Creep Life Using the European Creep Collaborative Committee (ECCC) Master Equation

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The experimental results are compared with the predicted creep life for standard solution annealed HR3C steel as determined using the ECCC published creep fracture master equation [50]:

tf = exp (β0 + β1σ0 + β2 σ0 2 + β3σ0 3 + β4 σ0 4 + β5 T) ,

(1)

where tf is the predicted time to fracture in hours, σ0 is the applied stress in MPa and T is temperature in K. The constant parameters take the values β0 = −26.7658463, β1 = −0.116574839, β2 = 0.00045779653, β3 = −9.44027818E−7, β4 = 6.22283614E−10 and β5 = 42,990.9766 [43]. The comparison of the experimental points with the calculated ones according the master equation is shown in Fig. 19. It can be seen that in the AR state the predicted and the experimentally determined times to fracture follow a one to one correlation at both testing temperatures. Whilst the simulated HAZ specimens also closely followed the same one-to-one correlation at 1023 K, they failed at 973 K significantly earlier than is the case for the AR material and the prediction. 5. Conclusions The short-term creep tests on austenitic stainless HR3C steel pipe in the AR state and after HAZ simulation by means of the GLEEBLE 3800 physical simulator were carried out at 923 and 1023 K and the applied tensile stress ranging from 100 to 350 MPa to evaluate the effect of welding on creep properties and behaviour. The main results are summarised as follows: 1. Following analysis of creep data, all the creep tests were performed in the region of power-law or dislocation creep. The rate controlling mechanism is most probably the climb of free intragranular dislocations. 2. The results show a clear detrimental influence of HAZ simulation on creep properties of HR3C steel at 923 K and very short creep exposures. However, with increasing time of the creep test, the effect of weld simulation becomes less significant. By contrast, no such deterioration of creep properties was detected at 1023 K. These findings agree with determined values of the weld strength factor (WSF). 3. The evaluated similar values of the stress exponent of the minimum

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