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P92 Weldment Issues OMMI (Vol. 1, Issue 2) August 2002

THE MECHANICAL PROPERTIES AND MICROSTRUCTURES OF 9% CHROMIUM STEEL P92 WELDMENTS

P.J. Ennis, Institut für Werkstoffe und Verfahren der Energietechnik IWV-2, Forschungszentrum Jülich GmbH, D-52425 Jülich’ Germany

Phil Ennis is at the Research Centre Jülich, Institute for Materials and Processes in Energy Systems, Julich, Germany. He is concerned withh materials development for advanced fossil-fired power plant, with emphasis on the creep rupture properties and microstructure of high chromium steels and nickel-based alloys. This paper is an updated and shortened version of the paper by P J Ennis and O Wachter published in VGB PowerTech, 1/98 pp 84-94

INTRODUCTION A new steel grade can only be considered for applications in piping systems and boiler construction when sufficient reliable information concerning the long-term behaviour is available. However, because welding is always used for joining, information on the basis steel alone is not enough; the behaviour of the weldments must be established in extensive and long-term tests for steels operating in the creep regime.

In this report the behaviour of P92 autogenous weldments will be discussed. The microstructure of the steel is of decisive importance for the creep rupture behaviour and therefore the relationships between structure and mechanical properties were investigated. Special emphasis was given to the creep rupture behaviour of the intercritical (softened) region in the heat affected zone (HAZ) of a weldment. In addition to stress rupture tests on actual weldments, weld simulation tests, in which the microstructure characteristics of the heat affected, softened zone were produced by an appropriate heat treatment, have been carried out using iso-stress tests; the results were compared with those obtained for the basic steel and for the complete weldment.

The susceptibility of P92 to temperature and strain induced embrittlement and to stress relief cracking was also investigated using weld simulation specimens; there is hardly any information in the literature on these effects.

MANUFACTURE OF THE WELDMENTS FOR TESTING Circumferential welds were made in a 300 mm diameter, 40 mm wall thickness pipe of commercial P92 manufactured by the Nippon Steel Corporation. The TIG welding wire Nittetsu YT-616 was used for the root pass and Nittetsu N-616 for the filler and surface weld layers. Both these weld metals were obtained from the Nippon Steel Welding Products and Engineering Company Limited. For the investigation of the mechanical properties of weldments, two welds designated were made by Mannesmann Anlagenbau. The welding conditions are given in Table 1. The two pipe welds were delivered in the as-welded condition and each pipe weld was cut into four

P92 Weldment Issues OMMI (Vol. 1, Issue 2) August 2002 2

segments, three of which were post-weld heat treated for 4 h at 710°C (low), 740°C (normal recommendation) and 770°C (high).

PREPARATION OF WELD SIMULATION SPECIMENS Weld simulation experiments were used in order to obtain test specimens with microstructures typical of the narrow intercritical region and the coarse grained region of the HAZ. The simulation of the very different over-heated and tempered microstructures in a weldment was accomplished using the Gleeble machine at the RWTH Aachen, Institute for Ferrous Metallurgy. With this apparatus, the thermal cycles and heating and cooling rates of a real weldment can be simulated.

For the simulation of a real weld thermal cycle, the characteristic value for the cooling is given by ∆t8/5, the time taken for the specimen to cool from 800°C to 500°C. To simulate the welding of a thick-walled pipe (40 mm) with a heat input of 14.4 kJ/cm, ∆t8/5 was set at 25 s. This value is somewhat conservative, calculations of the welding cycles given in [2] yielding ∆t8/5 of 12 s. The decision to set ∆t8/5 at 25 s was made considering the following factors. Firstly, a comparison with P91 would be possible [3] and secondly this value covers heat inputs up to about 25 J/cm, which are typical for manual arc welding of 20 mm wall thickness pipes in the overhand position with a pre-heat of at least 200K. This heat input may be regarded as representative for a ∆t8/5 value of 25 s [4], because for this weldment a ∆t8/5 of 24 - 30 s was determined using rod electrodes of 3.25 mm diameter. In addition, a heat input of 25 kJ/cm with a ∆t8/5 of 25 s is typical for submerged arc welding processes.

The weld simulation was carried out on 15 mm diameter bars machined from the pipe; test specimens were subsequently machined for short and long term stress rupture testing and for impact testing. Preliminary simulation tests had shown that for peak temperatures on the surface at the middle of the bar of 900, 1100 and 1300°C, and a heating rate of 150 K/s, a ∆t8/5 of 25 s could be achieved. Under these conditions, a useful specimen length of 20 - 30 mm with a homogeneous microstructure was obtained.

Investigation of the properties of the intercritical region was performed on specimens which had been given single-cycle Gleeble treatments with peak temperatures of 800, 850, 875, 900, 925 and 950°C. For evaluation of the properties of the coarse grained region of the HAZ which is of relevance for TIE, SIE and SRC, the Gleeble weld simulation treatments consisted of a single cycle with a peak temperature of 1300°C. For the second series of TIE specimens, a second cycle with a peak temperature of 900°C was applied after cooling to room temperature. In all the weld simulation experiments, there was no hold time at the peak temperature, the specimen being cooled immediately after reaching the peak temperature specified.

TEST PROGRAMME AND RESULTS The test programme for the weldments and the weld simulation specimens is summarized in Tables 2 and 3.


Weldments Chemical analysis of the weld deposit

The chemical compositions of the weld deposits of the weldments W1A and W1B made using Nippon Steel filler, together with the Nippon Steel specifications for the parent steel and weld deposit, are given in Table 4. The higher Mn (1.6%) and Ni (0.63%) contents of the weld deposit than for the parent steel (specified Mn and Ni contents of 0.6% max and 0.4% max respectively [5] should be noted, as should the very low S content of below 0.004% for the filler metal.

Macro- and microstructure

Figure 1 shows the macro- and microstructures of the weldment W1A, post-weld heat treated 740°C/4 h. The macrostructure across the weldment and the microstructures of the different regions in the weld metal and the heat affected zone (HAZ) are shown. The macrostructure is characterized by a very even, multi-layer weld deposit with a tempered HAZ region, recognizable as a 2 - 3 mm wide zone.

The microstructure of the weldment W1A revealed by TEM examination and shown in Figure 2 exhibited in the as-welded condition tempered martensite of typical lath form and many M23C6 particles precipitated on the prior austenite grain boundaries and subgrain boundaries. Occasionally islands of δ ferrite were observed which were surrounded by M23C6 particles. A determination of the transformation temperature of the weld deposit yielded an Ac1 temperature of 775°C. A post-weld heat treatment of 740°C/4 h caused substantial recovery effects to occur, but the martensitic structure with M23C6 precipitates on the grain boundaries remained clearly visible, Figure 2b. A few non-metallic inclusions of up to 1 µm diameter were seen.

In the fine grained (softened) HAZ of the weldment W1A in the as-welded condition, a weakly tempered martensitic structure was found. The martensite laths were not so well defined and the M23C6 particles were not evenly distributed, Figure 3a. In the post-weld heat treated condition, a subgrain structure, partially polygonized, was observed with a low dislocation density, Figure 3b.

Hardness

Hardness profiles (HV5) across the weldment were measured. Figure 4 shows the results obtained on the weldments, which were post-weld heat treated 740°C/4 h, for different positions in the wall of the pipe (inner, middle and outer regions of the wall).

The hardness profiles for the different positions have the same general characteristics, a maximum hardness in the weld metal and a steep decrease in the HAZ with a hardness trough. The levels of hardness in the weld metal and the width of the trough depended on the position in the pipe wall.

In the root region (inner region of pipe wall and the region of the most extensive tempering), the hardness trough is, if at all, difficult to recognize, the hardness reduction in the trough being less than 20 HV5 points below that of the parent material. In the wall middle and outer parts, the hardness trough was clearly observed, with a reduction of 20 to 30 HV5 points. In a hardness profile measurement at an angle of 70° to the fusion line to provide increased resolution of the hardness trough region, a reduction in hardness of more than 40 HV5 points was found; the hardness was 180 HV5 over a width of nearly 2 mm, as shown in Figure 5.


In W1A post-weld heat treated 740°C/4 h, the maximum hardness in the weld metal was between 290 and 330 HV5 and in W1B 330 - 350 HV5. The variation in the post-weld heat treatment temperature from 710 to 770°C did not lead to any marked difference in the hardness values measured at the pipe wall middle.

Impact strength

The weldment W1B was tested in the as-welded condition and after post-weld heat treatments of 4 h at 710, 740 and 770°C. The results are shown in Figure 6. The P92 matching filler metal of weldment W1B exhibited higher notch toughness after the post-weld heat treatments than in the as-welded condition; there was no appreciable effect of the heat treatment temperature. The plateau hardness was reached at a test temperature of about 100°C; the parent material exhibited the plateau at room temperature. The maximum impact energy of 150 J was a little below that of the parent material with 190 J.

The HAZ of the weldment W1B exhibited an increase in impact energy after heat treatment as was observed for the weld metal. Bearing in mind the large scatter in the toughness values for the parent material, the influence of post-weld heat treatment temperature can be regarded as only slight. The HAZ impact energy values were around the lower scatter band of the Av-T curve of the parent material and the hardness plateau value was even higher than that of the parent material.

Creep rupture strength

The results of the iso-stress rupture testing at 120 MPa of the weldment W1A are shown in Figure 7; the mean line for P92 was derived using the MCM method applied to all available P92 data [1]. In general, the stress rupture lives of the weldment in the as-welded condition and post-weld heat treated lay below those of the parent material. There was a tendency for the specimens post-weld heat treated 740°C/4 h to show the lowest strengths. It should be noted that two of the specimens tested at a stress of 120 MPa ruptured in the weld metal at times shorter by a factor of 6 at 612°C and 4 at 625°C than observed for the parent material. All other specimens ruptured in the HAZ. The rupture strengths of all W1A specimens, however, were above the lower scatter band (mean values less 20% in the stress direction) of P92 parent material.

Weld simulation Softened zone

- Softening temperature The determination of the weld simulation cycle temperature needed to produce a softened microstructure was carried out according to Table 3. The peak temperature in the Gleeble HAZ simulation cycle was varied between 800 and 950°C and the hardness HV10 measured in the as- simulated condition and after a heat treatment of 740°C/2 h. For the specimens heat treated 740°C/2 h, the tensile strength at 600°C and a strain rate of 10-5 s-1 was also determined.

In the as-simulated condition, a weak minimum in the hardness, 245 HV10, was seen at around 850°C. For peak temperatures above 875°C there was a sharp increase in hardness to values of around 410 HV10, which are typical for untempered martensite 8. In the heat treated condition the hardness level was about 50 HV10 points lower than in the as-simulated condition for peak temperatures up to 875°C; specimens subjected to simulation peak temperatures of 900 and 925°C


exhibited after additional heat treatment 740°C/2 h low hardnesses typical for tempered martensite.

The hot tensile tests at 600°C and a slow strain rate of 10-5 s-1, similar to a short time creep test, indicated a minimum tensile strength for a simulation peak temperature of 850°C followed by heat treatment 740°C/2 h, Figure 9.

- Stress rupture strength

The results of the stress rupture tests carried out at 612°C, 145 MPa on weld simulation specimens with different peak temperatures and heat treated 740°C/2 h also exhibited a minimum stress rupture life for a peak temperature of 850°C, as shown in Figure 10. In comparison with the results obtained for the basis steel in the as-received condition, the weld simulation specimens showed significantly shorter rupture times, which were of the order of the rupture lives exhibited by the parent material specimens that had been tempered at 835°C. For a more detailed iso-stress evaluation, a simulation peak temperature of 850°C was chosen on the basis of the above results. The specimens were heat treated 740°C/2 h after weld simulation.

The results of the extended iso-stress tests are shown in Figure 11, in which the rupture life is plotted as a function of test temperature for stresses of 120 and 145 MPa. The unfavourable values measured for the weld simulation specimens and for the parent material specimens tempered at 835°C compared with the strengths exhibited by parent material in the normal tempered condition (770°C) are clearly seen.

- Microstructure The microstructures of two weld simulation specimens were examined, namely a specimen heat treated 740°C/2 h after weld simulation at a peak temperature of 850°C and a specimen with the same weld simulation and heat treatment that had been creep rupture tested at 637°C/120 MPa, rupture life 725 h, rupture strain 52%.

The microstructure of the weld simulation specimen showed in comparison to the as-received condition clear evidence of polygonisation of the martensitic structure, Figure 12. The subgrain boundaries were relatively heavily decorated with M23C6 precipitates and the dislocation structure was inhomogeneous. In the subgrain interiors, the dislocation density was 3.3 ± 0.7 x 1014 m-2. The dislocations were generally associated with M(C,N) particles, Figure 12. After creep testing for only 725 h at 637°C, the martensite lath structure was no longer recognizable; the microstructure was characterized by large equiaxed subgrains. The dislocation density had decreased even more and M23C6 was found in the form of elongated particles associated with the Laves phase on the subgrain boundaries, Figure 12b.

Coarse grained zone - Temperature induced embrittlement (TIE) The notched impact strength of weld simulation specimens was measured and the results are shown in Figure 13. The post-weld heat treatment of 740°C/2 h led to an increase in the impact energy of the specimens given the double cycle simulation (Tmax 1300°C + 900°C) to more than 130 J from the 20 - 30 J measured in the as simulated condition. The improvement was not so marked for the single cycle simulation specimens. In the as-simulated double cycle condition, however, the impact energy remained at a low level of about 30 J up to test temperatures of


120°C. The single cycle simulation in contrast showed a slight increase from 20 - 60 J in the temperature range 20 to 120°C.

- Strain induced embrittlement (SIE) The influence of the strain imposed on the HAZ by the stress relief heat treatment at 740°C on the impact strength of the HAZ was determined. The results, which are shown in Figure 14, indicate that strains in the range 0.3 to 1.5% at 740°C did not have a detrimental effect on the impact energy at room temperature, measured using sub-size v-notch specimens. This is also the case if the results for the HAZ in the as-welded condition and for the parent material in the as-received condition are included after conversion of the impact energy to the sub-size specimens using the divisor 2.7 [1]. - Stress relief cracking (SRC) The planned tests to investigate the sensitivity of P92 weldments to stress relief cracking could not be successfully carried out. The temperature dependence of the rupture strain of coarse grained HAZ simulated specimens in low strain rate tensile tests at 550 to 775°C could not be determined because the specimens were fracturing outside the HAZ simulated region, the gauge length of the specimens. Even reducing the cross-section of the coarse grain simulated region by 50% did not induce fracture in the desired specimen region. Taking into account the results of the SIE tests described above, the SRC tests were discontinued.

DISCUSSION

Weldments Impact strength

From the results shown in Figure 6, in which the impact energies of weld deposit and HAZ of the weldments W1B are plotted as a function of temperature, there are two aspects to be considered, the impact strength of the weld deposit and that of the HAZ.

The weldment W1B with an impact energy of only 8 J at room temperature lies clearly below the specified minimum of 27 J for the average impact energy of the parent material as given in TRD 201 [6]. Values above 27 J were only attained for the weldment W1B at test temperatures of 50°C and above. At test temperatures of 100°C and above, impact energies of more than 150 J were obtained. If one considers that the impact strength decreases during creep loading, then some restrictions are needed to avoid brittle fracture in pipe pressurisation tests at low temperatures. As at the beginning of the P91 weld deposit development, there is for P92 further development necessary concerning impact strength improvement, unless it is permissible to carry out the pipe pressurisation test at 50°C or above. This would require design code approval.

Creep rupture strength

Specimens of the weldment W1A exhibited, with the exception of two specimens (612 and 625°C, 120 MPa) that ruptured in the weld deposit, the expected behaviour in creep rupture tests. That is, the rupture times more or less lay within the scatter band of P92 and the ruptures were transcrystalline in the HAZ. The reasons for the two ruptures in the weld deposit were not investigated further. The results of TEM investigation of the weld deposit and the associated hardness determinations did not provide any direct indications.


A comparative evaluation of the creep rupture behaviour of parent steel pipe in the as-received condition (AR), of the weldment W1A post-weld heat treated 740°C/2 h, of specimens given a weld simulation treatment with Tmax 850°C and subsequent heat treatment 740°C/2 h (WS3) and of the parent material after heat treatment 1070°C/2 h + 835°C/2 h (HT2) shows that the stress rupture life at 600 - 662°C and 120 MPa decreased in the order given above, see Figure 11. This is in agreement with the results of TEM investigations of the microstructure in the HAZ, which showed an increasing recovery of the microstructure and a continuous decrease in the mean dislocation density (Figure 3).

Weld Simulation Creep rupture behaviour

Based on the experience of the formation of the softened zone in the HAZ of weldments, preliminary tests were carried out and a weld simulation cycle with a peak temperature of 850°C was selected as a condition under which P92 would be sensitive to softening. This was revealed, not very clearly, in the hardness determinations after different weld simulation treatments (Figure 8) and in the slow strain rate tensile tests carried out at 600°C which showed a minimum tensile strength for the 850°C weld simulation, Figure ?. The reason for the loss in tensile strength for the 850°C weld simulation and the strength increase observed for weld simulations at 900°C and above is the same as given in [1] for the hardness loss and the loss of creep strength after a tempering treatment at 850°C. The carbon solubility in newly formed austenite at just above the Ac1 temperature of 825°C for the parent material is still so low that it leads to a decrease in the strength. With increasing temperature and increasing solubility of carbon in the austenite being formed, the strength increases to reach a maximum for temperatures above the Ac3 temperature, that is, for a fully martensitic microstructure. It is perhaps surprising that such effects can occur in very short times (in the range of some seconds) as indicated by the results of the iso-stress tests results.

The creep rupture behaviour of the 850°C weld simulation specimens compared to that of the 835°C tempered specimens conforms to the results of the TEM examinations. In the weld simulated condition WS3, the mean dislocation density was 3.3±0.7 x 1014 m-2 and after 835°C tempering (HT2) it was 2.3±0.6 x 1014 m-2. The microstructure after weld simulation, Figure 12 also indicated clear signs of recovery in the form of polygonisation in the martensitic structure, as for the 835°C tempered specimens. The specimen examined by TEM after creep testing (637°C, 120 MPa, 725 h, 52.3% rupture strain, condition WS3) confirmed the accelerated degradation of the microstructure; even after only 725 h test duration, there was no longer any indication of a martensitic structure.

From the above comments, a certain weakness of the steel P92 in the HAZ for long term service is recognized. How pronounced the weakening in the HAZ proves to be in practice in comparison with P91 can only be assessed when the long term testing of the weldments has been completed.

At this point, however, one thing should be mentioned. Tests reported in [7] concerning the sensitivity of weldments to Type IV cracking under creep loading (creep cracking in the HAZ, perpendicular to the highest stress direction in weldments stressed perpendicular to the fusion line) have shown that such cracks form preferentially and early when a marked softened zone in the HAZ is present, for example in P91 weldments and in contrast not in 12Cr1MoV weldments. Although in Europe longitudinally welded pipes, which would be especially sensitive to Type IV cracking, are not normally used for live steam components, careful attention should be paid during


the design stage to ensure that there is no high degree of multiaxiality in the principal stress in circumferential welds. The axial stresses, which are only half the circumferential stresses, should not be allowed to lead to a high degree of multiaxiality due to excessive bending. Based on the results obtained in the present work on weld simulation specimens, the above comments apply for P92.

Embrittlement caused by welding and post-weld heat treatment

The embrittlement caused by welding and subsequent heat treatment has been studied using coarse grain weld simulation experiments. The characteristic elements of the testing methods are tests to determine whether the steel exhibits a tendency for • temperature-induced embrittlement (TIE), • strain-induced embrittlement (SIE), • stress relaxation cracking (SRC).

The TIE and SIE tests have shown clearly that the coarse grained region of the HAZ in P92 weldments does not show any tendency for embrittlement after post-weld heat treatment. The impact energies determined in the TIE tests, Figure 13, confirm that the highest impact strengths are obtained when a coarse grain weld simulation at 1300°C is followed by a 900°C weld simulation cycle; a heat treatment 740°C/ 2 h after weld simulation produced a tempered martensitic structure. The results for the double weld simulation cycle 1300°C/900°C with no subsequent heat treatment allow the conclusion that an austenite of high hardness (470 HV5) may have formed during the 900°C cycle with the related uptake of carbon which led to the formation of a martensitic structure of low impact strength. The difference in the impact energies of the as weld simulated specimens and of weld simulated specimens given a heat treatment at 740°C means that in the HAZ of an actual weldment differences in impact strength are to be expected.

The SIE tests did not indicate any tendency for strain-induced embrittlement. The following metallurgical mechanism is proposed in [8, 9]. For low alloy fine grained construction steels, it is possible that during stress relieving heat treatments a marked strengthening of the grain interiors can be produced as a result of the precipitation of the microalloyed elements as fine dispersed particles. Glide processes for the relief of weld induced stresses have to operate mainly over grain boundaries. This results in a reduction in the deformability and is regarded as a preliminary step towards SIE. For P92 such differences in the precipitation processes and strengths of grain interiors and grain boundaries do not apparently occur, so that the tendency for SIE can be classed as low. This is in agreement with the experience obtained for 12Cr1MoV [10].

The fact that the SRC specimens always fractured outside the 1300°C weld simulation region with unusually high reduction of area justifies the assumption that in this region a high deformability must exist. Taking into account also the reduction of area values found in the normal creep tests, it becomes apparent that the ductility of the weld simulation specimens (rupture strains of over 50%) is significantly higher than the ductilities of the other specimens tested. It is therefore suggested that the rupture position of the 1300°C weld simulation specimen corresponds to the position at which the maximum temperature was only around 850°C and therefore the high deformation values were observed. More detailed investigation is required.

The SRC tests could not be carried out in the planned manner; to accomplish this, creep specimens that have the desired 1300°C weld simulation microstructure over the whole gauge length are required.


CONCLUSIONS In this work, the following aspects of P92 weldments with matching filler metals have been investigated:

- the influence of different post-weld heat treatments on the microstructural parameters important for the creep rupture strength;

- the creep behaviour of weldments, in comparison with that of the parent material;

- creep rupture behaviour of the softened HAZ determined using weld simulation specimens;

- HAZ embrittlement (coarse grained region) resulting from post-weld heat treatment, also investigated using weld simulation specimens.

The stress rupture strength of P92 decreased in the following order:

a) parent material

b) weldments ruptured in the HAZ

c) specimens with microstructure typical of the softened HAZ (weld simulation)

d) parent material tempered at 835°C instead of 775°C.

The P92 weldment made using matching filler (W1A) exhibited a creep rupture strength within the scatter band of the parent steel (stress mean value ± 20%). The rupture occurred generally in the HAZ and was transcrystalline.

Because of the likelihood of Type IV cracking in the HAZ under creep loading, it should be assured that the weldment is not subjected to the maximum stress across the weld or that a high degree of multiaxiality as a result of inappropriate design is not obtained in the HAZ.

The impact strength of the weld deposit of the matching filler weldment W1B was very low at room temperature; the impact energy was below 10 J. With increasing test temperature, however, the impact strength increased rapidly and at 100°C, impact energies above 150 J were obtained. Notwithstanding, some further optimisation appears to be needed. The temperature dependent toughness should be taken into account when conducting pressure testing of pipes.

Neither in the HAZ of an actual weldment nor in the HAZ coarse grain weld simulation (peak temperature 1300°C or 1300/900°C, no hold time, cooling from 800 to 500°C in 25 s, with or without subsequent heat treatment 740°C/2 h) was there a tendency for temperature or strain induced embrittlement found. The susceptibility for stress relief cracking was also considered to be low.

REFERENCES [1] Wachter, O. and Ennis, P. J., Investigation of the Properties of the 9% Chromium Steel 9Cr0.5Mo1.8WVNb with a View to its Application for Pipework and Boilers Operating at Elevated Temperatures, Part 1: Microstructure and Mechanical Properties of the Basis Steel, VGB Kraftwerkstechnik 77 (1977), Number 9, 699-713


[2] Buchmayer, B., WEZ-Kalkulator, ein Softwarepaket zum Berechnen schweißtechnischer Kennwerte, Schweißen und Schneiden 41 (1989), H.2, 69-75 [3] Brühl, F., Verhalten des 9% Chromstahles X10 CrMoVNb 91 und seiner Schweißverbindungen im Kurz- und Langzeitversuch, Dissertation (May 1989); Technische Universität Graz [4] Aberer, R., Schweißverhalten des 9% Chromstahles P91, Diplomarbeit, TU Graz, Institut für Werkstoffkunde, Festigkeitslehre und Materialprüfung, Abteilung Werkstoffe und Schweißtechnik, Sept.1988 [5] Nippon Steel Corporation, Data Package for P92 Ferritic Steel (9Cr-0.5Mo-1.8W-Nb-V), Jan. 1993, Second Edition March 1994 [6] Technische Regeln für Dampfkessel, TRD 201 (Herstellung): Schweißen von Bauteilen aus Stahl, Fertigung - Prüfung, (Juni 1989), TRD 201, Anlage 1 (Herstellung):Schweißen von Bauteilen aus Stahl, Richtlinien für die Verfahrensprüfung (Juni 1989), Carl Heymanns Verlag KG und Beuth Verlag GmbH (1994) [7] Coussement, C., Verelst, L., Multiaxial Creep Behaviour of Welded Components in High Strength Ferritic/Martensitic Creep Resistant Steel, 5th International Conference on Materials for Advanced Power Engineering 1994, 3-6 Oct 1994, Liège, Belgium, Kluwer Academic Publishers [8] Kussmaul, K. und Ewald, J., Assessing Heat Affected Zone Properties of Pressure Vessel Weldments Using Simulation Techniques, International Journal of Pressure Vessel and Piping, Volume 5, No. 2 (1977), 85-101 [9] Kussmaul, K. und Ewald, J., Assessment of Toughness and Cracking in the Heat-Affected Zone of Light Water Reactor Components, Third International Conference on Pressure Vessel Technology- Part II: Materials and Fabrication (1977), published by The American Society of Mechanical Engineers [10] Kalwa, G. und Schnabel, E., Umwandlungsverhalten und Wärmebehandlung der martensitischen Stähle mit 9 bis 12 % Chrom, VGB-Konferenz "Werkstoffe und Schweißtechnik im Kraftwerk 1989", Essen


welding plan for pipe welds of P92

test material

weldments W1A and W1B (pipe, 300 x 40 mm)

weld preparation J-form with 8 to 10° angle root step height 2 mm

welding method root

TIG/ Argon 99.99%

filler and capping Manual metal arc filler metal: root pass Nippon wire Nittetsu YT616

d = 2 mm filler and capping Nippon rod electrode

Nittetsu N 616 d = 3.25 mm

pre-heat temperature > 200°C, generally 235 to 245°C inter-pass temperature < 250°C welding position underhand, pipe axis horizontal welding current TIG: ca.135 A SMA ca.120 A welding voltage ca. 24 V welding speed ca. 12 cm/min oscillation width 2-3 x core wire diameter, mm

layer thickness < 2 mm heat input 14.4 kJ/cm ∆t8/5 12 s hydrogen effusion heat treatment >250°C, 2 h post-weld heat treatment each weldment:

1 quarter - as welded 1 quarter - 710°C/4 h 1 quarter - 740°C/4 h 1 quarter - 770°C/4 h heating rate: 150K/h to 500°C; 80K/h from 500°C cooling rate: 80K/h to 500°C, 150K/h to RT

Table 1: Welding parameters for the P92 weldments


weldments W1A/W1B weld deposit and HAZ

iso-stress rupture tests (all heat treatments)

across the weldment W1A heat

treatment condition +

OM W1A W1B

hardness W1A W1B

NIS

W1B

TEM

W1A

temp °C

stress MPa

expected rupture time*, h

test type

OM

600 12 000 MSM x

612 3 000 MSM x

625 145 1 000 SSM x

637 300 SSM x

AW x x x x 650 100 SSM x

P1 x x x 662 30 SSM x

P2 x x x x 600 100 000 MSM x

P3 x x x 612 30 000 MSM x

625 8 000 MSM x

637 120 2 500 SSM x

650 700 SSM x

662 250 SSM x

675 70 SSM x * rupture time data for basis steel, as received (1070°C/2 h+ 775°C/2 h) + AW: as welded; P1 post-weld heat treated 710°C/4 h (low); P2 740°C/4 h (normal); P3 770°C/4 h (high) SSM: single specimen machine MSM: multi-specimen machine NIS: notched impact strength as function of test temperature; 1 ISO v-notch specimen for each temperature: notch location: for HAZ, 0.5±0.3mm from the fusion line, crack propagation in the welding direction; for weld deposit, perpendicular to surface, middle of weld deposit, crack propagation in the welding direction Hardness: HV5 measurements perpendicular to fusion line for indication of softened region in the HAZ OM: optical metallography TEM: transmission electron microscopy

Table 2: Test matrix for investigation of weldments W1A and W1B, parent material, matching P92 filler metal, with various post-weld heat treatments


objective Gleeble cycle, Tmax ,°C, without holdtime

test

1. Determination of the properties of the softened HAZ

a) determination of Tmax for softening

• 800, 825, 850, 875, 900, 925 with/without 740°C/4 h • 800, 825, 850, 875, 900, 925 with 740°C/2 h

• hardness HV10 • HTT at 600°C, strain rate 10-5 s-1

800, 825, 850, 875, 900 with 740°C/2 h (conditions WS1 to 5)

iso-stress tests*: • 145 MPa; 612, 625°C

b) creep rupture test of softened microstructure

Tmax for softening selected 850°C# with heat treatment 740°C/2 (condition WS3)

iso-stress tests • 120 MPa: 600, 612, 625, 637#,**, 650, 662°C • 145 MPa: 612, 625°C

2. Investigation of the behaviour of the coarse grained zone

a) stress relief cracking (SRC) 1300 • HTT at 550, 600, 650, 700, 775°C; strain rate 3.10-4 s-

1 • HTT at temperature of minimum fracture strain, strain rate 10-6 s-1 **

b) temperature induced embrittlement (TIE)

1300, 1300+900 Gleeble treatment with/without 740°C/2 h

impact energy / temperature curves, 1 ISO v-notch specimen per temperature

c) strain induced embrittlement (SIE)

1300 NIS at 20°C on specimens tested in HTT at 740°C and removed at strains of 0.3, 0.7, 1.0, 1.5%

HTT: hot tensile test NIS: notched impact strength, ISO v-notch, sub-size specimens * expected rupture time for as received condition, test types as in Table 8 **OM # TEM WS: weld simulation with subsequent heat treatment

Table 3: Test matrix for weld simulation tests on P92


element

weldment Nippon weld filler metal **

specification ** parent material

wt % W1A W1B SAW * P92 C 0.085 0.082 0.086 0.06 - 0.13 Si 0.33 0.17 0.36 <0.50 Mn 1.60 1.60 1.53 0.30 - 0.60 P nd nd 0.006 <0.020 S 0.001 0.002 0.003 <0.010 Cr 8.5 8.55 8.85 8.00 - 9.00 Mo 0.54 0.54 0.49 0.30 - 0.60 W 1.60 1.65 1.62 1.50 - 2.20 V 0.24 0.25 0.30 0.15 - 0.25 Nb 0.086 0.085 0.050 0.03 - 0.10 B <0.008 <0.008 nd <0.006 N 0.038 0.037 0.041 0.030 - 0.070 Ni 0.63 0.62 0.58 <0.40 Al <0.03 <0.03 nd <0.04

nd: not determined * SAW (shielded arc welding), with rod electrodes ** data of Nippon Steel [5]

Table 4: Chemical compositions of welds W1A and W1B (filler metal similar to P92)


Figure 1: Macro- and microstructure of the weldment W1A (P92 matching filler), post-weld heat treated 740°C/4 h

HAZ Weld deposit

Coarse and fine grain HAZ Fine grain HAZ


(a) (b)

Figure 2: TEM thin foils taken from the weld deposit of weldment W1A; a) as welded, b) after post-weld heat treatment 740°C/4 h

(a) (b) Figure 3: TEM thin foils of the softened HAZ of weldment W1A; a) as welded, b) post-weld heat treated 740°C/4 h


150

175

200

225

250

275

300

325

350

0 2 4 6 8 10 12 14 16 18 20

distance, mm

Hv5

mid-wall

inner

outer

fusion lines

P92 weldmentPWHT 4 h/740°C

Figure 4: Hardness profiles across the weld after post-weld heat treatment 740°C/4 h

150

170

190

210

230

250

270

290

310

330

350

0 1 2 3 4 5

distance, mm

hard

ness

, Hv1

0

fusion line

weld deposit

softened zone

parent material

Figure 5: Hardness profile at a 70° angle to the fusion line for the weldment W1A (P92 filler), after post-weld heat treatment 740°C/4 h, showing location of specimens for TEM investigations of microstructure


0

50

100

150

200

-120 -80 -40 0 40 80 120

temperature, °C

impa

ct e

nerg

y C

VN, J

scatter-band for parent material

as welded

4h/710°C

4h/740°C

4h/770°C

post-weld heat treated

notch in weld deposit

0

50

100

150

200

250

-120 -80 -40 0 40 80

temperature, °C

impa

ct e

nerg

y CV

N, J

as w elded

4h/710°C

4h/740°C

4h/770°C

scatter-band for parent material

notch in HAZ

Figure 6: Influence of tempering treatment on the impact energy of a) the weld deposit (WD) and b) the heat affected zone (HAZ) of weldment W1B (filler P92); values for parent metal given for comparison


10

100

1000

10000

100000

600 625 650 675test temperature, °C

time

to ru

ptur

e, h

as welded4h/710°C4h/740°C4h/770°C835°C temper

iso-stress rupture tests120 MPa

mean ine for P92

Figure 7: Rupture time from iso-stress tests at 120 MPa on the weldment W1A (filler P92) as welded and after post-weld heat treatments; the mean line for the parent material and the values for P92 tempered at 835°C are given for comparison

150

200

250

300

350

400

450

500

800 850 900 950

T max weld simulation, °C

hard

ness

, Hv1

0

no heat treatment

heat treated 2h/740°C

Figure 8: Hardness as a function of peak temperature in weld simulation; specimens with and without heat treatment after simulation


200

300

400

800 850 900 950

T max w eld sim ulation, °C

tens

ile s

treng

th, M

PaP 92 weld simualtiontensile test at 600°Cstrain rate 10 - 5 s - 1

Figure 9: Tensile strength at 600°C determined in tensile tests at low strain rate as a function of peak temperature Tmax in the weld simulation; specimens were heat treated 740°C/2 h after weld simulation

0

1000

2000

3000

4000

5000

6000

Parent775°C

temper

Parent835°C

tem per

Tm 800°C Tm 825°C Tm 850°C Tm 875°C Tm 900°C

weld simulation specimens

time

to ru

ptur

e, h

612°C 145 MPa

Figure 10: Time to rupture for weld simulation specimens with different peak temperatures Tm; data for the parent steel and the parent steel tempered at 835°C are provided for comparison


10

100

1000

10000

100000

575 600 625 650 675

test temperature, °C

time

to ru

ptur

e, °C

iso-stress tests 145 MPa

10

100

1000

10000

100000

575 600 625 650 675

test temperature, °C

time

to ru

ptur

e, h

parent 775°C temperparent 835°C temperweldment PWHT 740°Cweld simulation HAZ

iso-stress tests at 120 MPa

Figure 11: Rupture time in iso-stress tests of weld simulation specimens, heat treated 740°C/2 h, compared with parent steel, as received and after tempering at 835°C:


Figure 12: TEM thin foils of weld simulation specimens with Tmax 850°C followed by heat treatment 740°C/2 h (WS3): a) before and b) after creep testing at 637°C, 120 MPa, 725 h, 52.3% rupture strain

0

20

40

60

80

100

120

140

1300°C 1300 P WHT 1300+900 1300+900 P WHT

coarse-grained weld simulation specimens

impa

ct e

nerg

y, J

20°C

40°C 80°C

Figure 13: Influence of the coarse grain weld simulation (single cycle Tmax 1300°C and double cycle Tmax 1300°C/900°C) on the impact strength for investigation of temperature induced embrittlement following welding and post-weld heat treatment


0

10

20

30

40

50

60

70

80

90

-0.25 0 0.25 0.5 0.75 1 1.25 1.5 1.75

tensile pre-strain in SIE specimens at 740°C, %

RT

impa

ct e

nerg

y, J

parent steel

weldment HAZ

TIE specimen

SIE

Figure 14: Influence of pre-strain in a tensile test at the tempering temperature (740°C) on the impact strength of coarse grain weld simulation microstructure (Tmax 1300°C) [to investigate susceptibility to strain induced cracking; comparison values for the parent material, for weldment W1B (post-weld heat treated 740°C/4 h) and for a TIE specimen (weld simulation Tmax 1300°C, heat treated 740°C/2 h)]

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