Microstructural Study of Weld Joint Made of Cast P91 Steel after Creep Testing

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1 Microstructural Study of Weld Joint Made of Cast P91 Steel after Creep Testing D. Jandová, J. Kasl and V. Kanta ŠKODA VÝZKUM s.r.o., Tylova 57, CZ , Pilsen, Czech Republic Abstract Trial weld joint was prepared using the GTAW SMAW method. Creep testing was carried out at temperatures ranging from 525 C to 625 C and stresses from 50 MPa to 240 MPa. Creep strength of weld joint was evaluated according the Larson-Miller parametric equation. At temperatures up to 575 C it falls into the ±20% scatter band of the creep strength of the base material; at higher temperatures it decreases bellow the bottom of this scatter band. Microstructural analysis, which included light, scanning and transmission electron microscopy, elucidated causes of creep failure. The failure was concentrated in the fine grain heat affected zone, in the region with a large area of prior austenitic grain boundaries, high density of coarse precipitates and relatively low density of fine vanadium carbonitrides. Fractures generally occurred in this region. Only one specimen exposed at 625 C for 29,312 hrs ruptured in weld metal. After this exposure many coarse particles of Laves phase were detected and remarkable growth of subgrains was observed in the weld metal. In addition, some particles of Z-phase were identified that caused partial dissolution of vanadium carbonitrides. Coarsening of particles of secondary phases and recovery of dislocation substructure resulted in material softening, plastic deformation and formation of cavities in the critical regions. Introduction The creep-resistant ferritic steels alloyed with chromium and molybdenum are widely used for production of components of boilers and turbine parts in steam power plants. In modern power plants the steam conditions exceed 600 C and 26 MPa (Ultra Super Critical Conditions). Material that underwent these severe conditions have to be stable for several decades; the service life of turbines would TABLE 1: CHEMICAL COMPOSITION OF THE BASE MATERIAL AND CONSUMABLES (WT. %) be at least hrs [1]. Nowadays even hrs are required for the life time of power plants. Grade P91 is 9Cr-1Mo-V steel, which is currently used for production of components of high efficiency power plants that operate at temperatures up to 585 C. On the base of this steel new grades were developed, that are designed for applications in USC conditions [2]. A lot of creep data was collected of P91 steel [3], [4]. The high creep strength of this steel is a result of solid solution hardening of Mo atoms, carbide-stabilized substructure hardening (caused by M 23C 6 carbide) and precipitation hardening of vanadium nitrides or carbonitrides [5], [6]. Welding is a common technology of joining the machine components in power plants. Welds have heterogeneous structure and are often susceptible to fracture. Failure is usually concentrated in a specific part of weld joint as a result of different kinetics of microstructural processes in individual regions like weld metal, heat affected zone and the base material. Nowadays the processes taking place during longterm creep exposures are often simulated using numerical models [7], however each model represents simplification of the real conditions and therefore it is always useful to investigate microstructure using specimens of real weld joints. A detailed study of precipitation and recovery of dislocation substructure can elucidate causes of preferential development of creep failure in some regions of weld joint [8], [9]. The paper deals with a study of trial weld joint that was produced by SKODA POWER a.s. within a frame of development of welding technology, that is used for production of steam piping and cast components of turbines. Some results of investigation of this weld joint and also of other similar and dissimilar weld joints of P91 steel have been published [10]-[13]. C Mn Si Cr Mo V Ni Nb N P S P Thermanit MTS Chromo 9V

Experimental procedures Weld joint of cast plates made of 9Cr-1Mo-V steel (P91 steel according ASME, GX12CrMoVNbN 91 steel in the Europe Standardization) was fabricated using GTAW & SMAW method.

Inductive heating with thermal isolation ensured a preheating temperature in the range (200-250) C. The interpass temperature was kept below 300 C.

The post-weld heat treatment (PWHT) was carried out at a temperature in the range from 740 C to 750 C for 2.5 hrs.

2 Experimental procedures Weld joint of cast plates made of 9Cr-1Mo-V steel (P91 steel according ASME, GX12CrMoVNbN 91 steel in the Europe Standardization) was fabricated using GTAW & SMAW method. The plates with dimensions of 500 x 150 x 25 mm in conditions after austenitizatization at 1050 C for 1.5 hrs, oil quenching and tempering at 750 C for 3.5 hrs were joined in PA position. Inductive heating with thermal isolation ensured a preheating temperature in the range ( ) C. The interpass temperature was kept below 300 C. Root pass was formed with Thermanit MTS3 and Chromo 9V was used as an electrode. The chemical composition of the base material and the consumables is given in Table 1. The post-weld heat treatment (PWHT) was carried out at a temperature in the range from 740 C to 750 C for 2.5 hrs. Integrity and mechanical properties of weld joint have been evaluated according to the welding standards EN 288-2,3. Required mechanical properties as well as results of nondestructive testing in accordance with quoted standards were satisfied. Smooth cross-weld specimens with a length of 92mm and a diameter of 8 mm were undergone creep testing. Fractographic and metallographic analyses of ruptured specimens were undertaken using light (LM) and scanning electron microscopy (SEM). Metallographic samples were prepared in longitudinal section of the creep test specimens. Vickers hardness measurement and metallographic analysis were performed. Macro and also microstructure was revealed using Villela Bain s reagent. Extraction carbon replicas and thin foils for transmission electron microscopy (TEM) were prepared from selected regions of weld joint; from the weld metal (WM), the coarse prior austenitic grain heat affected zone (CG HAZ), the fine prior austenitic grain heat affected zone (FG HAZ) and the base material (BM). Foils were produced using jet electropolishing in 6% solution of perchloric acid in methanol at a temperature of -40 C. Particles of secondary phases and substructure were observed using microcope JEOL JEM 1200EX and EDX microanalyser Oxford Instruments INCA 300. The bottom line of the ±20% scatter band of the creep rupture of the base material, which is usually permitted for weld joints, is represented with a dash grey line. Vertical dash lines correspond to Larson-Miller parameter for time to rupture of hrs and temperatures of 550 C, 575 C and 600 C. The creep strength of the weld joint falls into the considered scatter band up to 575 C. At higher temperatures it rapidly decreases. At 625 C and 60 MPa the creep strength of the weld joint is 27 % lower than the creep strength of the base material. Stress [MPa C P = T (25 + log t) 575 C 600 C Figure 1. The creep rupture strength of the weld joint in comparison with the creep rupture strength of the base material GX12CrMoVNbN 9 1 steel. Creep tested specimens ruptured mostly in the heat affected zone. Only one exception has been registered up to this time; specimen exposed at 625 C/ 50 MPa /29,312 hrs fractured in the central part of the weld metal (Fig. 2). Elongation was usually a few percents. Only specimen exposed at the lowest temperature and the highest stress and also specimen exposed at 625 C/50MPa/29,312hrs revealed considerable macroplastic deformation (elongation about 10%). Fracture surfaces were covered with oxide layers; therefore it was difficult to distinguish the fracture mechanism. Fast fracture was restricted to small marginal regions (Fig. 3). Results Ceep-testing was carried out for 16 specimens at temperatures in the range from 525 C to 625 C and stresses from 40 MPa up to 240 MPa. Time to the rupture ranges from 1,061 hrs to 33,189 hrs and two tests are still running (marked with arrows in Fig. 1). Creep rupture strength was evaluated using the Larson-Miller parametric equation P = T [C + log τ], where T represents temperature given in degree Kelvin, C is a specific constant for a given material (C = 25) and τ means time to fracture in hours. The creep rupture strength of crossweld specimens (a black line in Fig. 1) was compared with the creep rupture strength of the base material (a grey line) a b Figure 2. Macrostructure of cros-weld creep tested specimens: a) 600 C/80MPa/25,818hrs - sample C10, b) 625 C/50MPa/29,312hrs - sample C14.

$Central parts of fracture surfaces indicated a dimpling morphology. At two specimens, which were exposed at 575 C, also intercrystalline facets were detected [11], [13].$ Some cavities that formed at the prior austenite grain boundary are indicated with arrows in Fig. 4. tests at conditions of 600 C/80MPa/25,818hrs and 625 C/60MPa/19,210hrs.

Some cavities that formed at the prior austenite grain boundary are indicated with arrows in Fig. 4. tests at conditions of 600 C/80MPa/25,818hrs and 625 C/60MPa/19,210hrs.

$Fast fracture 5 mm Figure 3. Fracture of the cross-weld specimen tested at 600 C/80MPa/25,818hrs sample C10. Figure 5.$ $Cavities were always observed near the fracture surface (Fig. 5). After creep tests at temperatures from 525 C to 575 C for approximately 20,000 hrs cavities were concentrated in FG HAZ.$ During long-term exposures at temperature 550 C and stresses 160 and 180 MPa (about 30,000 hrs) separate cavities appeared in WM and BM, however most of them were again located in FG HAZ.

During long-term exposures at temperature 550 C and stresses 160 and 180 MPa (about 30,000 hrs) separate cavities appeared in WM and BM, however most of them were again located in FG HAZ.

3 Central parts of fracture surfaces indicated a dimpling morphology. At two specimens, which were exposed at 575 C, also intercrystalline facets were detected [11], [13]. It can be concluded that all fractures occurred as a result of growth and coalescence of cavities, that arose in inner parts of the tested bar. Some cavities that formed at the prior austenite grain boundary are indicated with arrows in Fig. 4. tests at conditions of 600 C/80MPa/25,818hrs and 625 C/60MPa/19,210hrs. After creep test at 625 C/50MPa/29,312hrs a lot of cavities were found in WM and FG HAZ (Fig. 6). Cavities were observed only exceptionally in the base material outside HAZ. Fast fracture 5 mm Figure 3. Fracture of the cross-weld specimen tested at 600 C/80MPa/25,818hrs sample C10. Figure 5. Cavitation failure in FG HAZ after the creep test at 575 C/120MPa/19,289hrs - sample C8. SEM micrograph of the etched metallographic sample. 10 m Figure 4. Fracture surface of sample C14 after the creep test at 625 C/50MPa/29,312hrs. SEM micrograph. Metallographic analysis also revealed cavitation failure. Cavities were always observed near the fracture surface (Fig. 5). After creep tests at temperatures from 525 C to 575 C for approximately 20,000 hrs cavities were concentrated in FG HAZ. During long-term exposures at temperature 550 C and stresses 160 and 180 MPa (about 30,000 hrs) separate cavities appeared in WM and BM, however most of them were again located in FG HAZ. Similar distribution of cavities was observed after creep Figure 6. Cavitation failure in the weld metal after the creep test at 625 C/50MPa/29,312hrs - sample C14. LM micrograph. Vickers hardness profile across the weld joint was determined for the weld joint before and after creep testing. Before creep testing hardness of the base material was 210 HV10 and the weld metal 226 HV10 (sample C). During creep tests up to 600 C hardness slightly increases in both the weld metal and the base material (sample C10), while after tests at 625 C with 19,210 and 29,312 hrs to the rupture (sample C14), hardness of the weld metal

$HV 10 260 240 220 200 180 C C10 C14 exposure at 550 C/160MPa/33,189hrs (Fig. 9). Coarse particles of Laves phase were identified in the weld metal using electron diffraction and EDX microanalysis.$ This phase can be distinguished beyond all doubts according high contents of iron and molybdenum and lower content of chromium.

This phase can be distinguished beyond all doubts according high contents of iron and molybdenum and lower content of chromium.

4 decreased to 203 and 181 VH10, respectively. Local maxima in CG HAZ near the fusion line and local minima in FG HAZ were detected (Fig. 7). HV C C10 C14 exposure at 550 C/160MPa/33,189hrs (Fig. 9). Coarse particles of Laves phase were identified in the weld metal using electron diffraction and EDX microanalysis. This phase can be distinguished beyond all doubts according high contents of iron and molybdenum and lower content of chromium. On the contrary high content of chromium and small amounts of vanadium, iron and molybdenum were detected in M 23C 6 carbides. Changes in density of vanadium carbonitrides were not found out BM WM BM Distance [mm] Figure 7. Hardness profile across weld joint: Before creep testing (sample C), after creep test at 600 C/80MPa/25,818hrs (sample C10) and after creep test at 625 C/50MPa/29,3128hrs (sample C14). Microstructure of the weld joint was tempered martensite. Dendrite segregation was evident in the base material. After post-weld heat treatment substructures of the base material and the weld metal were similar. Martensitic laths were divided into subgrains, relatively coarse chromium rich particles of M 23C 6 occurred at grain and subgrain boundaries. In some of laths fine intragranular cubic V(C,N) precipitate was observed (Fig. 8). In addition, fine orthorombic (Cr,V) 2C precipitate was identified in the weld metal. Density of coarse particles was higher in the base material than in the weld metal and distribution of fine intragranular precipitates was irregular. Coarse carbides were rarely observed in CG HAZ, while density of fine precipitate was relatively high. On the contrary a lot of coarse particles occurred in FC HAZ and fine vanadium carbonitrides were observed only exceptionally. Dislocation density was approximately m -2 in the base material and in CG HAZ and m -2 in the weld metal. Microstructure in FG HAZ consisted of relatively small grains with subgrains almost free of dislocations. During creep exposures recovery occurred, coarse particles slightly grew, (Cr,V) 2C carbide particles dissolved and new vanadium nitrides precipitated. Dissolution of precipitates after creep exposures up to 550 C resulted in a hardness decrease of the weld metal, while precipitation of V(C,N) particles after exposures at 600 C resulted in a hardness increase. After creep test at 575 C/140MPa/10,031hrs any particles of M 2C precipitate were not observed; size and density of other secondary phases and also dislocation density were the same as before creep testing. Other substructural changes revealed after long-term 200 nm Figure 8. Substructure of the base material before creep testing. TEM micrograph. 500 nm Figure 9. Substructure of the weld metal after the creep test at 550 C/160MPa/33,189hrs sample C5. TEM micrograph. Significant changes of substructure, which resulted in a remarkable hardness decrease, were observed in the weld metal after exposure at 625 C/50MPa/29,312hrs. A size of

subgrains increased in comparison to dislocation substructure at conditions before creep testing and also to long-term creep exposure at 550 C (Fig. 9 and 10).

Laves phase particles were also detected in the base material, however they were not by a large size. Laves phase often nucleated at the surface of oxide particles. hardening became lower.

$Rows of diffraction patterns were detected that corresponded to (001) planes of tetragonal Z-phase with interplanar distance of 0.739 nm [14],[15].$

5 subgrains increased in comparison to dislocation substructure at conditions before creep testing and also to long-term creep exposure at 550 C (Fig. 9 and 10). Particles of Fe2Mo Laves phase became coarser. Their size often exceeded 1 m in the weld metal (Fig. 11). Laves phase particles were also detected in the base material, however they were not by a large size. Laves phase often nucleated at the surface of oxide particles. hardening became lower. In the weld metal and in the base material some plate-like particles of Z-phase were found. Rows of diffraction patterns were detected that corresponded to (001) planes of tetragonal Z-phase with interplanar distance of nm [14],[15]. One particle was observed, which generated diffraction patterns of both crystallographic variants of Z-phase tetragonal and also cubic face centered one (Fig. 9). Similar particles that revealed both above mentioned crystallographic symmetries have been observed in 12CrMoVNb steel [16], [17]. Z-phase 500 nm Figure 10. Substructure of the weld metal after the creep test at 625 C/50MPa/29,312hrs sample C14. TEM micrograph. Oxide 200 nm Figure 12. A particle of Z-phase in the weld metal after the creep test at 625 C/50MPa/29,312hrs sample C14. TEM micrograph. Discussion Laves phase 200 nm Figure 11. A cluster of particle in the weld metal after the creep test at 625 C/50MPa/29,312hrs sample C14. TEM micrograph. A growth of Laves phase resulted in depletion of solid solution of molybdenum and decrease in substitution hardening. Simultaneously a density of fine vanadium carbonitrides decreased and consequently also precipitation Critical zones from a point of view of creep failure in the weld joint depended on conditions of creep testing. After creep exposures at temperatures from 525 C to 600 C and short-term exposures at 625 C fractures occurred in the fine prior austenite grain heat affected zones, in the regions with relatively high density of coarse M23C6 carbides distributed at grain/subgrains boundaries and relatively low density of fine intragranular vanadium carbonitrides. On the other hand during long-term exposure at 625 C creep failure was concentrated in the weld metal. After exposure at 625 C/50MPa/29,312hrs massive particles of Fe2Mo Laves phase besides of M23C6 carbides and V(C,N) precipitate were observed in the weld metal. Vanadium carbonitrides were detected only exceptionally. In addition some thin plate-like particles of Z-phase were detected. Precipitation processes taking place during creep exposures resulted in different hardening and consequently different level of plastic deformation and cavitation failure in individual zones of weld joint.

6 It is well known that strain of the steel investigated occurs by dislocation creep at high temperatures. Deformation of individual grains increases if grains are soft (without fine precipitate and dislocation). The steps are formed at grain boundaries that together with coarse secondary phases serve as nucleation centres for cavities. Growth of cavities is promoted by high speed grain boundary diffusion. During the creep exposure dislocation density in dislocation walls increases and the subgrain boundaries act as channels of a high diffusion like the grain boundaries. Large surface of grain/subgrain boundaries promotes a growth of cavities, their coalescence and the crack propagation, therefore the cavitation failure occurs in FG HAZ predominately. Combination of dislocation and cavitation creep results in a dimpling morphology of fracture surfaces. During creep exposures at temperatures above 600 C for approximately 30,000 hrs coarsening of Laves phase was indicated especially in the weld metal. Precipitates of Laves phase nucleated on the surface of other particles, especially of oxides, that are always present in the weld metal as a result of welding procedure. Occurrence of massive Laves phase particles and decrease in the solid solution hardening were the main reasons of fracture in the weld metal. Precipitation of Z- phase, which caused partial dissolution of fine vanadium carbonitride, was of secondary importance, because fine plate-like particles of Z-phase sparsely occurred in both the weld metal and also the base material. Conclusions The creep strength of the weld joint examined falls into desirable ±20% scatter band of creep strength of the base material up to 575 C. At higher temperatures the creep strength of the weld joint does not reach desirable creep strength in comparison to the base material. Critical zones from a point of view of creep failure at temperatures up to 600 C are fine grain heat affected zones. At higher temperatures coarsening of Fe 2Mo Laves phase causes a solid solution softening and the crack propagation in the weld metal. Some particles of Z-phase were detected in the weld metal and also in the base material after creep test at 625 C for about 30,000 hrs. This phase causes dissolution of vanadium carbonitride and decrease in the creep strength during longterm creep exposures. Acknowledgements This work was supported by Grant project MSM and 1P05OC024 COST 536 from the Ministry of Education, Youth and Sports of the Czech Republic. References [1] J. Hald, "Creep resistant 9-12%Cr steels-long-term testing, microstructure, stability and development potentials," in Proc. Superhigh strength steels, Milano: Associazione Italiana di Metalurgia, 2005 [CD-ROM] 146. [2] M.E. Staubli, K.H. Mayer, T.U. Kern, R.W. Vastone, "COST 501/COST 522, The European collaboration in advanced steam turbine materials for ultra efficient, low emission steam power plants," in Proc. Parsons 2000, Advanced Materials for 21 th Century Turbines and Power Plants, A. Strang et al. Eds., London: IOM, 2000, p.98. [3] H.K.D.H. Bhadeshia, "Design of ferritic creep resistant steels, ISIJ International, vol. 41, pp [4] A. Strang, V. Foldyna, A. Jakobová, Y. Kuboň, V. Vodárek, J. Lenert, "Factors affecting the prediction of long term creep rupture properties of microstructurally unstable 9-12%Cr power plants steels," in Proc. Advancees in Turbine Materials, Design and Manufacturing, A. Strang Eds., IOM, Newcastle upon Tyne 1997, p [5] I. Cipolla, J. Gabrel, "New creep rupture assessment of grade 91", in Proc. Super-high strength steels, Milano: Associazione Italiana di Metalurgia, 2005 [CD-ROM] 162. [6] G. Eggeler, "The effect of long-term creep on particle coarsening in tempered martensite fertritic steels," Acta Metallurgica, Vol. 37, 1989, p [7] S. Concari, "WELDOM overview of industrial research and experience related to welds and data collection in the WELDOM databank," in Proc. Mechanics and Material in Design, J.S. Silva Gomes and S.A. Meguid EDs., Porto: Universidada do Porto & INEGO, 2006 [CD- ROM] A [8] R.W. Vastone, "Alloy design and microstructural control for improvement of 9-12%Cr power plant steels", in Proc. Materials for Advanced Power Engineering 2002, Vol. II, J. Lecomte-Beckers et al. EDs, Liege: Forschungszentrum Jülich, 2002, p [9] G. Moscal, A. Hernas, J. Pasternak, "Structural stability and properties of weld joints of new creep-resistant steel grades with 9 or 12% Cr contents applied in operation of power generation sector over 30,000hrs", in Proc. Materials for Advanced Power Engineering 2006, Vol. III, J. Lecomte-Beckers et al. EDs, Liege: Forschungszentrum Jülich, 2006, p [10] D. Jandová, J. Kasl, V. Kanta, "Creep resistance of similar and dissimilar weld joints of steel P91", Materials at High Temperatures, Vol. 23, No.3-4,, 2006, pp [11] D. Jandová, J. Kasl, V. Kanta, E. Folková,"Weldment of cast steel P91 - creep testing and microstructure evaluation, " in Proc. Mechanics and Material in Design, J.S. Silva Gomes and S.A. Meguid EDs., Porto: Universidada do Porto & INEGO, 2006 [CD-ROM] A [12] D. Jandová, J. Kasl, E. Folková, V. Kanta,"Microstructural studies of similar and dissimilar weld of P91," in Proc. Materials for Advanced Power Engineering 2006,. J. Lecomte-Becker et al. EDs., Lieg;e: Forschungszentrum Jülich GmbH 2006, p [13] D. Jandová, J. Kasl, V. Kanta, " Long-term creep testing and microstructure evaluation of P91 steel weld joints ", in Proc. BALTICA VII Life management and Maintenance for Power Plants, Helsinki: VTT Technical Research Centre of Finland and Finishing maintenance Society, 2007, pp [14] A. Strang, V. Vodárek, "Z-phase formation in martensitic 12CrMoVNb steel," materials Science and Technology, Vol. 12, 1996, pp [15] V. Vodárek, H.K. Danielsen, F.G. Grumsen, J. Hald, A. Strang, "Ellectron diffraction studies on (Nb,V)CrN particles in 12CrMoVNbN steels," in Proc. Materials for Advanced Power Engineering 2006, Vol. III, J. Lecomte-Becker et al. EDs., Liege: Forschungszentrum Jülich GmbH 2006, p [16] H.K. Danielsen, J. Hald, "Behavior of Z-phase in 9-102% Cr steels," Energy Materials, Vol. 1, No. 1, 2006, pp [17] H.K. Danielsen, J. Hald, "Z-phase in 9-12 Cr steels observations and thermodynamic modeling," in Proc. Materials for Advanced Power Engineering 2006, Vol. III, J. Lecomte-Becker et al. EDs., Liege: Forschungszentrum Jülich GmbH 2006, p.1275.

7 Published in: Welding in the World 52 (2008) pp (Special Issue Proceedings of the IIW International Conference Safety and Reliability of Welded Components in Energy and Processing Industry, July 2008, Graz, Austria. P. Mayr, G. Posch, H. Cerjak Eds. Graz University of Technology, Graz).

CAUSES OF FAILURE OF WELD JOINT DURING LONG TIME CREEP TESTING. Dagmar JANDOVÁ, Josef KASL, Eva CHVOSTOVÁ

Abstract CAUSES OF FAILURE OF WELD JOINT DURING LONG TIME CREEP TESTING Dagmar JANDOVÁ, Josef KASL, Eva CHVOSTOVÁ a ŠKODA VÝZKUM s.r.o., Tylova 1/57, 316 00 Plzeň, Česká republika, damar.jandova@skodavyzkum.cz