Heat affected zone microstructures in supermartensitic stainless steels

P Woollin* and D Carrouge**
* TWI Ltd, Cambridge
** University of Cambridge, Department of Materials Science & Metallurgy

Paper presented at conference on Supermartensitic stainless steels, Brussels, Belgium, 3 - 4 October 2002

Abstract

The use of supermartensitic stainless steels has expanded rapidly in recent years but characterisation of weld heat affected zone (HAZ) microstructure has been the subject of only limited study. Consequently, the currentunderstanding of HAZs is limited, particularly with respect to corrosion, stress corrosion cracking and sulphide stress cracking behaviour. The paper describes the microstructures in the various regions of a supermartensitic stainlesssteel HAZ and considers effects of re-heating from subsequent passes and PWHT.

1. Introduction

Supermartensitic stainless steels offer a combination of high strength, corrosion resistance and weldability that is particularly attractive for oil and gas flowlines. Compared to more traditional nickel-alloyed, soft martensiticstainless steels with around 13%Cr4%Ni, they have a lower carbon level (about 0.01% compared to ≥0.03%), an addition of molybdenum to improve corrosion resistance (typically 2%) and an increase of nickel content (typically to 5-6.5%). Leansupermartensitic grades, with little or no molybdenum and less nickel, are also available, as are intermediate grades. ^[1] Some grades also have a small titanium addition, to limit secondary hardening. Typical compositions are shown in Table 1.

Table 1 Chemical analyses of the range of steels examined

There is a substantial body of data indicating that crack-free welds with good HAZ toughness can be made without need for pre-heat and postweld heat treatment. ^[2,3] Hence, the steels are widely considered to have good 'weldability'. However, a number of HAZ corrosion and stress corrosion issues remain ^[4-6] and there is a need for improved understanding of the HAZ microstructure development in order to address these. Unfortunately, the HAZ microstructure is very fine scale and all relevant features cannot be fully resolved under anoptical light microscope, which has hampered understanding of HAZ behaviour.

This paper reviews the current understanding of supermartensitic stainless steel HAZ microstructure and serves as an update to the paper presented by Gooch et al at the 1999 Supermartensitic Stainless Steel Conference. ^[7]

2. Parent steel microstructure

The parent steels are supplied in the tempered condition. Whilst precise details are proprietary, the steels typically contain several percent of stable revert austenite together with tempered martensite containing a population ofcarbides. Published literature indicates that it is possible to attain 20-30% stable austenite in 13%Cr steel with 4-6%Ni, if tempering is performed at about 100°C above Ac ₁. ^[7,8] Both chromium and molybdenum may be incorporated into carbides during tempering, together with minor elements such as V and Ti. For conventional nickel alloyed martensitic stainless steels, M ₂C and M ₇C ₃ (where M is a metal atom) form in the range 400 to 450°C, and contribute to secondary hardening, whilst M ₂₃C ₆ is dominant at above 500°C. M ₂C is stabilised by the presence of Mo, thus Mo alloying encourages secondary hardening on tempering. ^[8,9] It has been reported that tempering at 620-680°C of the coarse grained HAZ of a weld in supermartensitic stainless steel with 2.7%Mo predominantly gives molybdenum carbides. ^[10]

3. The various regions of a supermartensitic HAZ

3.1 General comments

The HAZ may be subdivided into several regions with different microstructures. ^[6,11,12] The sequence of transformation on heating is as follows. Below Ac ₁ some tempering, i.e. carbide precipitation, may occur. Above Ac ₁, martensite transforms to austenite and the proportion of austenite increases until the steel is fully austenitic at Ac ₃. Martensite remaining in this temperature interval will be tempered. There is then a single-phase austenite range until Ac ₄, when delta ferrite formation begins. The microstructure becomes fully ferritic at Ac ₅ and remains single phase until the melting point. On cooling, the resulting microstructure reflects the peak HAZ temperature attained and five distinct HAZ regions can be identified:

(i) that which was fully delta ferrite (>Ac ₅, which is >around 1370-1420°C, Table 2),
(ii) that which was a mixture of delta ferrite and austenite (Ac ₄-Ac ₅, which is around 1330-1370°C to 1350-1420°C),
(iii) that which was fully austenite (Ac ₃-Ac ₄, which is around 720-800°C to 1330-1370°C),
(iv) that which was a mixture of austenite and martensite (Ac ₁-Ac ₃, which is around 540-680°C to 720-800°C) and
(v) that which did not exceed Ac ₁ but which exceeded the minimum temperature for carbide precipitation, i.e. around 400°C.

Table 2 Transformation temperatures for three supermartensitic stainless steels, measured with a heating rate of 10°C/min

The characteristics of the various regions are illustrated in Figs.1-4 and discussed below. A typical microhardness distribution for a single pass TIG weld HAZ in 12Cr6Ni2Mo steel is shown in Fig.5.

3.2 HAZ that was fully delta ferrite at peak temperature

There are few published data available relating to Ac ₄ and Ac ₅ temperatures for supermartensitic stainless steels, although some values calculated using MTDATA have been reported, ^[12] indicating temperatures in the range 1240-1280°C for Ac ₄ and 1270-1300°C for Ac ₅ for a lean grade (11Cr1.5Ni) and a high alloy grade (12Cr6.5Ni2.5Mo) respectively. Subsequent measurements for the same steels have indicated that actual temperatures are somewhat higher than predicted, being1330-1370°C for Ac ₄ and 1350-1420°C for Ac ₅, Table 2, for a heating rate of 10°C/s. ^[11] These temperatures are very high and approaching the melting point. Consequently, the fully ferritic region is narrow, typically 100-200 microns for conventional TIG/MIG welding processes, Fig.1. This region exhibits grain growth due to rapid diffusion in the ferrite phase and, on subsequent cooling, the austenite retains a coarser grain size than the parent. However, identifying prior austenite and prior deltaferrite grain structures is difficult by optical light microscopy. Consequently, this region may also be referred to as the grain coarsened HAZ. On cooling, the austenite growth into the delta ferrite is predominantly Widmanstatten innature with some allotriomorphic and intragranular austenite, ^[11,12] Fig.2 and 3. Austenite formation on cooling is apparently controlled by diffusion of substitutional elements as filaments of delta ferrite are retained between prior-austenite units to room temperature in the higher alloygrades and microanalysis indicates enrichment of the delta ferrite in Cr and Mo and depletion of Ni, Table 3. ^[11,12] Delta ferrite retention is probably a result of the low rate of back diffusion of substitutional elements into austenite during the transformation from ferrite to austenite on cooling, when compared to the opposite reaction onheating. Steels with higher levels of molybdenum and nickel retain higher volume fractions of delta ferrite than the lean alloys. The volume fraction of retained delta ferrite is small in actual HAZs, typically well below 10%, althoughpreliminary studies have shown higher levels (exceeding 10% delta ferrite) in furnace heat treated samples and indicated a weak link to the cooling rate ^[11] Additional work is required to establish the effect of welding parameters and the resulting thermal cycle on HAZ microstructure. On cooling, ferrite is retained preferentially in pre-existing areas of segregation in the steel.Enerhaug et al reported partial melting in this region of the HAZ, which is presumably related to pre-existing segregation also. ^[5]

Table 3 EDX microanalysis (wt%) of delta ferrite in weld HAZ of 12Cr6.5Ni2.5Mo steel and 11Cr1.5Ni steel. Measurements were made in region that was mixed delta ferrite and austenite at peak temperature but similar analyses apply tothe grain coarsened region that was fully delta ferrite at peak temperature.

Little, if any, austenite is retained to room temperature after heating to in excess of Ac 5, so that the final microstructure consists of martensite with a small fraction of retained delta ferrite. M _s is of the order of 250°C for a 12Cr6Ni2Mo steel. ^[1]

3.3 HAZ that was mixed delta ferrite and austenite at peak temperature

The region of HAZ that was partially ferritic at peak temperature does not show any evidence of appreciable prior austenite grain coarsening, presumably indicating that the austenite remaining at peak temperature prevents deltaferrite grain growth. Due to the small temperature range between Ac ₄ and Ac ₅, this region is narrow, of the order of 50-100 microns wide for a typical TIG/MIG weld, Fig.2. Austenite reformation on cooling is apparently not Widmanstatten in nature in this area but identifying the growth morphology is difficult due to the fineness of the microstructure in this region. However, on coolingdelta ferrite is once again retained to room temperature at a level that is typically less than 10%. In this case, the delta ferrite morphology is apparently intergranular, presumably with respect to prior austenite grains andenrichment of Cr and Mo and depletion of Ni is observed once more, Table 3. ^[11]

3.4 HAZ that was fully austenitic at peak temperature

The part of the HAZ that was fully austenitic at peak temperature may revert on cooling to a structure consisting entirely of virgin martensite or a mixture of martensite and stable austenite (remaining from the original parentsteel microstructure). Untempered, virgin martensite and retained austenite structures cannot be distinguished readily under an optical light microscope although colour-etching techniques can be useful. Depending on the extent ofdiffusion, previously stable austenite may remain stable or become unstable on cooling. The time at temperature required for sufficient diffusion to convert original stable austenite into unstable austenite (and hence martensite oncooling) has not been established.

This region typically shows a range of hardness and includes the hardest region of the HAZ, higher than that in the regions that are fully or partially delta ferrite at peak temperature, Fig 5. The primary factor controlling hardness of martensite is the carbon content (or carbon plus nitrogen content), with prior austenite grain size as a secondary factor. ^[13] Hardness increases with increasing carbon plus nitrogen and decreasing prior austenite grain size. Any delta ferrite will tend to lower hardness. Likely explanations for the hardness peak observed therefore include the fineraustenite grain size in this region compared to the grain coarsened HAZ and the absence of ferrite. A further possible explanation is carbon and nitrogen enrichment of some of the martensite, as a consequence of partitioning of carbonand nitrogen between stable revert austenite and unstable austenite during heating. The last possibility is supported by the observation of highest hardness in simulated HAZ microstructures when some austenite is present rather than infully martensitic structures.

3.5 HAZ region that was mixed austenite and tempered martensite at peak temperature

This region shows a mixture of dark etching grains, typically appearing darker than those of the parent steel presumably due to additional tempering and light etching grains of virgin martensite, Fig.4. There may also be some stable, revert austenite formed during heating. When the peak temperature exceeds Ac ₁ by a fairly small margin, partial transformation to austenite occurs. For furnace heat treatments of conventional nickel alloyed martensitic stainless steels, the austenite formed just above Ac ₁ is sufficiently enriched in austenite stabilising elements (e.g. C and N in particular) that its M _s temperature is reduced to below room temperature and it remains stable on cooling. ^[7-9] The temperature range over which this occurs is from Ac ₁ to around Ac ₁ + 100°C. ^[7-9] At higher temperatures, the amount of austenite formed on heating increases, the enrichment of the austenite is consequently less pronounced and hence the fraction of stable, revert austenite decreases. For conventional nickelalloyed martensitic stainless steel, the stable, revert austenite decreases to zero for heating to around 200°C above Ac ₁. ^[7-9] There are no published data available relating to formation of revert austenite during a weld thermal cycle in supermartensitic stainless steel. However, it is expected that the kinetics of transformation will be such that anequilibrium level of revert austenite will not form during a typical welding cycle.

The hardness of this region shows a transition from the parent steel value to the higher transformed HAZ levels, reflecting the variable mixture of tempered and untempered martensite, Fig.5. Softening in this region compared to parent steel, by about 10 VPN, has been reported. ^[6] This indicates that the additional tempering of untransformed martensite may outweigh the effect of formation of some virgin martensite but the extent of any such effect will presumably depend on the original tempering treatmentof the parent steel.

Fig.5. Microhardness variation in single pass TIG bead-in-groove weld HAZ in 12Cr6Ni2Mo steel, made with superduplex filler wire. Arc energy = 0.35kJ/mm. Approximate locations of isotherms corresponding to the various phase transformation temperatures are shown

3.6 HAZ that did not exceed Ac ₁ at peak temperature

There will also be a degree of HAZ tempering below the Ac ₁ temperature, especially for the lean grades (Ac ₁ around 650-680°C), but this will be of a limited nature for the higher alloyed grades due to the low Ac ₁ temperature of these materials (about 550°C, Table 2). Again the extent of any softening will depend on the original tempering condition of the steel. A neural network for prediction of Ac ₁ of supermartensitic stainless steels as a function of heating rate has been developed. ^[11] This is available from the following site: www.msm.cam.ac.uk/map/neural

4. Reheating effects of a second thermal cycle

A second weld thermal cycle can have a number of effects on HAZ microstructure. When reheated above Ac ₁, a degree of grain refinement is observed in the grain coarsened region (i.e. the region that exceeded Ac ₅ during the first cycle) due to nucleation and growth of new austenite grains. In other parts of the HAZ exceeding Ac ₁, reformation of austenite will be expected, to a first approximation, to give a resultant microstructure comparable with that associated with a single thermal cycle. However, where a non-equilibrium level ofstable austenite exists after one cycle, the second thermal cycle will increase the time available for diffusion and may reduce the level of such austenite.

A second cycle also allows an opportunity for carbide precipitation above and below Ac ₁ (in the approximate range 400°C-Ac ₃), whereas the low M _s temperatures of these steels will effectively prevent precipitation within any virgin martensite formed during the first weld cycle. This may be very significant with respect to intergranular corrosion of weldHAZs, which to date has been reported only in weld roots, i.e. areas subjected to multiple thermal cycles. ^[4,5] Cr ₂₃C ₆ precipitation on prior austenite boundaries occurs at above about 450-500°C, ^[8,9] so any sensitisation might be expected to develop in regions heated to between 450°C and the upper limit of Cr ₂₃C ₆ formation without adequate chromium mobility to prevent formation of depleted zones. Published information suggests the upper temperature for sensitisation is around 550-600°C, ^[14,17] although some effect of bulk composition is anticipated.

5. Postweld heat treatment

The effect of conventional PWHT on mechanical properties, in terms of softening through carbide precipitation and revert austenite formation is well established for nickel-alloyed 13%Cr steels with ≥0.03%C ^[7,9] but little specific information is available for the supermartensitic grades. ^[10] Maximum softening of conventional martensitic stainless steels, in general, is associated with Cr ₂₃C ₆ precipitation at above 500°C and for nickel free steels, tempering is typically performed at 700°C. For nickel alloyed supermartensitic grades, the low Ac ₁ temperature necessitates a somewhat lower tempering temperature. For tempering between 620 and 680°C a supermartensitic steel with 2.7%Mo was found to form predominantly molybdenum carbides, ^[10] which tend to give secondary hardening. Consequently, the Mo-alloyed supermartensitic steels are fairly resistant to softening on tempering. For 13%Cr4%Ni steels, tempering at about 600°C has been adopted but forsupermartensitic steels with around 6%Ni and 2%Mo a brief treatment at 650°C has been used when PWHT was considered necessary. ^[3,6] This is considerably in excess of Ac ₁ according to measurements made at 10°C/min, ^[10] although it should be noted that Ac ₁ is heating rate dependent and generally increases with heating rate. ^[7] This brief 650°C PWHT was reported to increase resistance to sulphide stress cracking, ^[6] although the mechanism of improvement was not identified. In this respect, it may be noted that Enerhaug et al concluded that behaviour of HAZs in mildly sour environments was controlled by the surface oxide rather than theunderlying microstructures. ^[6] However, in test specimens without surface oxide present, Hara and Asahi found that the presence of delta ferrite increased susceptibility to sulphide stress cracking. ^[18] The effect was attributed to precipitation of chromium carbides/nitrides at ferrite/martensite boundaries. Also, in laboratory studies, it has been reported that brief PWHT at 650°C increases resistance to intergranularcorrosion, suggesting that healing of chromium-depleted zones may occur even in a very short time. ^[5] The diffusion distance for chromium is of the order of only 10nm for 5 minutes at 650°C, using the published diffusion coefficient for Cr in Fe ^[19] but the extent of any Cr depleted zone has not been established yet.

In order to be able to control reliably the mechanical and corrosion properties of supermartensitic weld HAZs, there is a need for additional data to establish the detailed effects of PWHT, in addition to the weld thermal cycle, onthe microstructure and surface oxide in the HAZ. The microstructure and surface oxide will vary with location within the HAZ and between the variety of supermartensitic steels currently available.

6. Concluding remarks

Since the 1999 Supermartensitic Stainless Steel Conference, the industrial exploitation of these steels has increased substantially. In parallel, some efforts have been made to develop understanding of weld HAZ microstructuredevelopment and effects of PWHT. Nevertheless, there is still limited quantitative data available concerning the precise details of delta ferrite, austenite, martensite and carbide formation for the variety of steels available, whichalthough of similar composition, may have significant differences in HAZ microstructure and hence properties.

Due to the lack of data on HAZ microstructure, it is very difficult, at present, to link the various features of a specific weld HAZ to the properties of relevance to service and to set meaningful quantitative limits, e.g. for weldprocedure qualification purposes. To date, some efforts have been made to limit weld area hardness but there is an absence of data on the effect of hardness on properties relevant to service and it has not been demonstrated thathardness is a controlling parameter. Currently, interest is focussed on corrosion and stress corrosion properties of the weld HAZ but attention is also required in relation to mechanical integrity issues and in particular the effectsof hydrogen. At present, there is no alternative to qualification of welding procedures for supermartensitic stainless steels via testing which simulates conservatively the likely installation and service conditions or by directcomparison with equivalent or less onerous service experience with the same grade.

7. Acknowledgements

The authors thank colleagues at TWI and University of Cambridge for assistance in preparing the paper and are particularly indebted to the late Trevor Gooch for his guidance and encouragement and to the late Alan Haynes, NiDI.

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