Linear friction welding of AISI 316L stainless steel

Materials Science and Engineering A 528 (2010) 680–690

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Linear friction welding of AISI 316L stainless steel Imran Bhamji a,∗ , Michael Preuss a , Philip L. Threadgill b , Richard J. Moat a , Adrian C. Addison c , Matthew J. Peel d a

Manchester Materials Science Centre, University of Manchester, Grosvenor Street, M1 7HS, UK Formerly with TWI Ltd., Cambridge, UK (now retired) c TWI Ltd., Cambridge, UK d University of Bristol, Queens Building, University Walk, Bristol BS8 1TR, UK b

a r t i c l e

i n f o

Article history: Received 14 April 2010 Received in revised form 14 September 2010 Accepted 15 September 2010

Keywords: Linear friction welding AISI 316L Delta ferrite Texture

a b s t r a c t Linear friction welding is a solid state joining process established as a niche technology for the joining of aeroengine bladed disks. However, the process is not limited to this application, and therefore the feasibility of joining a common engineering austenitic steel, AISI 316L, has been explored. It was found that mechanically sound linear friction welds could be produced in 316L, with tensile properties in most welds exceeding those of the parent material. The mechanical properties of the welds were also found to be insensitive to relatively large changes in welding parameters. Texture was investigated in one weld using high energy synchrotron X-ray diffraction. Results showed a strong {1 1 1} 1 1 2  type texture at the centre of the weld, which is a typical shear texture in face centre cubic materials. Variations in welding parameters were seen to have a significant impact on the microstructures of welds. This was particularly evident in the variation of the fraction of delta ferrite, in the thermo-mechanically affected zone of the welds, with different process parameters. Analysis of the variation in delta ferrite, with different welding parameters, has produced some interesting insights into heat generation and dissipation during the process. It is hoped that a greater understanding of the process could help to make the parameter optimisation process, when welding 316L as well as other materials, more efficient. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Linear friction welding (LFW) is a solid state joining process, which is established as a niche technology for the fabrication of integrally bladed disk (blisk) assemblies in aeroengines [1,2]. The process involves a part, which is reciprocating in a linear manner (i.e. back and forth on one axis), being forced against a stationary part (Fig. 1). The friction between the parts creates heat, which together with the applied force causes a plasticised layer to form at the interface. Much of this plasticised material, in the weld region, is removed from the weld because of the acting forces, and the expelled material forms the so-called flash. This results in a loss of length of the overall component. Towards the end of the process the two parts are effectively forged together with some plasticised material remaining at the weld line. Practically all published literature on the subject has concentrated on the joining of materials that are used for aeroengine applications (i.e. titanium and nickel alloys, e.g. [3,4]). However, LFW is not restricted to joining these materials, although their low thermal conductivities certainly help to maintain sufficient heat at

∗ Corresponding author. E-mail address: [email protected] (I. Bhamji). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.09.043

the interface. Consequently, the possibility of LFW the austenitic stainless steel AISI 316L has been explored. AISI 316L is an engineering stainless steel widely used in a variety of industries and applications, including applications in the nuclear industry [5]. The primary aim of this paper was to assess the LFW process for producing parts in 316L stainless steel. A further aim was to identify the effect of welding parameters on microstructural features in the weld region, and so develop optimised welding parameters when LFW 316L. Austenitic stainless steels are readily welded by a number of different welding processes, and welds with good mechanical properties have been produced with the more conventional arc welding processes (e.g. [6]) as well as power beam process, such as laser (e.g. [7,8]) and electron beam welding (e.g. [9]). However, when welding austenitic stainless steels these fusion welding processes produce harmful hexavalent chromium fumes [10], which means solid state welding processes, where fumes are not produced, may be beneficial in some circumstances. Welding of austenitic stainless steels with the solid state friction stir welding process has already been demonstrated (e.g. [11,12]). The present paper concerns the LFW of such materials. As all common welding methods can be used to join austenitic stainless steels, process selection is a complex issue and will be determined by the intricacies of the particular application. How-

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Fig. 1. Schematic diagram of the linear friction welding process.

ever, there are a number of advantages to using friction welding, over other processes, which could make it viable for certain purposes. Key among these is the fact that the processes are solid state and therefore solidification problems (e.g. hot cracking, porosity, segregation, etc.) are avoided. Solidification problems have been encountered in arc welding of austenitic stainless steel (e.g. [13]), and are also an issue in laser welding [14]. However, good welds can be produced with all of these welding processes if adequate care is taken over the process development; of course all welding processes require care over their development to produce satisfactory welds. Friction welding also provides parts with low distortion, and the reliability and repeatability of the parts is increased as the processes are fully automated. The major disadvantage of friction welding processes is the high capital cost of the equipment, and this is especially the case for LFW, although the capital costs are reducing as the design of the machines develop. This high cost is likely to mean that the LFW processes will not attain widespread use for joining stainless steels in the short term, although it could be viable for certain high value added niche applications. Process selection between the friction welding family of processes is simpler as the different processes are normally used for different part geometries. Friction stir welding is generally used for butt welds in thin sheet, rotary and inertia friction welding are suitable for tubes and cylinders, whilst LFW is usually used for non axi-symmetric blocks. Although the LFW process brings about many advantages over fusion welding processes, a particular feature of the process when welding austenitic stainless steels (␥ phase, face centre cubic crystal structure) may be the formation of delta (␦)-ferrite and sigma (␴)-phase in the heat affected or thermo-mechanically affected zones of the weld. The high temperature ␦-ferrite phase, which has a body centred cubic (BCC) crystal structure, forms in fusion welds of 300 series austenitic stainless steels (for example [15]). The phase forms in regions heated to above about 1150 ◦ C [5], with some of it retained on cooling. In fusion welds a small amount can actually be beneficial as it limits the occurrence of hot cracking. However, when large fractions of ␦-ferrite are retained (more than

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about 12% [16]) it tends to reduce corrosion resistance and toughness of the weld region [15]. It is also thought that it can rapidly decompose during cooling to the ␴ and ␥ phases in friction stir welds, below about 1000 ◦ C, as strain and recrystallisation during the process can significantly accelerate the kinetics of the phase transformation [17]. As the LFW process also imparts significant plasticity on the welded samples the ␦ → ␴ + ␥ transformation may also take place in this process. The intermetallic ␴ phase is detrimental to weld properties because of its low fracture toughness and its tendency to deplete the adjacent austenite of chromium, thus making the material susceptible to corrosion [18]. Although the decomposition of ␦ to ␴ is not usually an issue in fusion welds, as the amount of process related plasticity and recrystallisation is small, it may still occur if the fusion weld is exposed to high temperatures for prolonged periods in service. As both large increases in ␦-ferrite and the formation of ␴ phase can be detrimental to weld properties, it is preferred that no ␴ phase is produced as a result of the LFW process and the amount of ␦ferrite within the microstructure is kept at a similar level to that in the parent material. Consequently an attempt has been made to optimise welding parameters for minimising the generation of ␦-ferrite. As ␦-ferrite is a high temperature phase, the welding temperature will have a large impact on the amount of the phase left in the microstructure after welding. There have been a few studies into the effects of welding parameters, during the LFW of different titanium alloys, on welding temperature [19–21]. However, the results from these studies appear to be contradictory, especially with regard to the contribution of applied pressure. Results from a study by Wanjara and Jahazi [21], on Ti–6Al–4V (Ti-64), suggest that as the applied pressure is increased a higher temperature is generated in the weld region. This was evident from an increasing prior ␤ grain size with increasing welding pressure. This was contradicted by Attallah et al. [19] who observed the prior ␤ grain size (in Ti–6Al–2Sn–4Zr–6Mo welds) becomes smaller with increasing welding pressure, concluding that cooler welds are produced when the welding pressure rises. This conclusion is supported by Romero et al. [20], who demonstrated that Ti-64 linear friction welds with low residual stresses can be produced if high welding pressures are applied. Heat has a major influence on residual stress development and low stress welds are indicative of a low welding temperature. There are a number of possible reasons for the discrepancies between different studies. These include differences in the welded material, parameters used (full disclosure of parameters was not given in [19,20]), sample geometry and aspects to do with the particular LFW machine (factors such as the efficiency of the machine in transferring power to the interface and the rigidity of part clamping will also have an effect on results). However more work is needed to identify the exact cause of the discrepancies and the actual relationship between welding parameters and welding temperature. There has been a relatively small amount of research into crystallographic texture development in linear friction welds, even though the material near the weld line is exposed to severe deformation during the joining process. To date textures have been studied in Ti-64 linear friction welds, and a very strong transverse ␣ texture, {1 0 1¯ 0}1 1 2¯ 0, and a {1 1 1} fibre ␤ texture has been reported close to the weld line [22]. Since variation in textures in 316L are thought to influence the materials fatigue properties [23] it is important to establish if strong texture development takes place during 316L LFW. As there is no publicly available information on textures in austenitic stainless steel linear friction welds it is difficult to predict the type of textures that would be expected. However, there has been some work on rolling textures in these materials and textures

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I. Bhamji et al. / Materials Science and Engineering A 528 (2010) 680–690 Table 1 Minimum, mid and maximum welding parameter values for 316L stainless steel linear friction welds. The same value (3 mm) was used for mid and maximum burnoff. Parameters for all min values: 25 Hz frequency, 1 mm amplitude, 60MPa friction pressure and 2 mm burn-off distance. Parameters for all maximum values: 45 Hz frequency, 2.5 mm amplitude, 200 MPa friction pressure, 2 mm burn-off distance.

Fig. 2. Schematic diagram of the parameter traces that are obtained during the linear friction welding process.

in linear friction welds may be similar. When rolling 316L different texture components were prevalent depending on the deformation conditions [24]. During cold rolling up to 40% reduction copper type textures (high intensities along the ␤ fibre, {1 1 2} 1 1 1  to {1 1 0} 1 1 2 ), together with a Goss component ({0 1 1} 1 0 0 ) were dominant. After larger reduction (which is more similar to LFW) a typical brass type texture was present (high intensities at {1 1 0} 1 1 2  and {0 1 1} 1 0 0 ), together with an increasingly prominent {1 1 1} fibre texture (although this component was still very minor relative to the others). 2. Experimental 2.1. Materials The material used for producing small scale linear friction welds in AISI 316L (Cr 17.3, Ni 9.8, Mo 2, Mn 1.9, Si 0.5, Fe balance, all in wt%; determined from experimental analysis of the base material) was supplied as 20 mm × 20 mm square drawn bar, which was saw cut into 75 mm lengths to produce weld samples. 2.2. Process parameters Samples were welded on an electromechanical LFW machine (a rotating crankshaft is attached to two cranks which provide the linear reciprocation; an alternative would be to use hydraulic systems to generate reciprocation) at TWI Ltd, Cambridge, UK. Preparation of the sawn ends, which would form the faying surfaces (welding faces), was limited to degreasing with acetone just prior to welding. There are a number of key parameters involved in LFW [2] (Fig. 2), which can have a significant impact on the joint properties. These parameters include:

Variable

Maximum

Mid

Min

Frequency (Hz) Amplitude (±mm) Friction pressure (MPa) Burn-off distance (mm)

45 3 240 3

40 2.5 160

25 1 60 1

at zero amplitude • forge pressure: the pressure, which is usually higher than the pressure during the frictional phase, that is applied at the end of the process (after oscillation has stopped), for a set period of time, to consolidate the joint • burn-off distance: the distance that the parts are allowed to shorten (shortening is caused by the removal of plasticised material from the weld) before the amplitude is decayed and the forge pressure is applied In this study four parameters (friction pressure, frequency, amplitude and burn-off distance) were varied to produce 22 different parameter combinations and 66 welds (3 welds at each parameter setting). In each weld a forge pressure that was the same as the friction pressure was applied for a period of 5 s. For each of the parameters that were varied, a minimum, mid and maximum value was determined from experience at TWI of friction welding materials similar to 316L. Each parameter was varied between its minimum and maximum value whilst keeping all other parameters constant at the mid value (8 different force settings, 4 frequency settings, 4 amplitude settings and 3 burn-off settings were used, in all cases whilst keeping other, non-variable, parameters at the mid level). In addition, welds were fabricated with all values at the minimum, mid and maximum values (giving 3 additional parameter settings). However, it was not always possible to use all of the extreme parameters in combination due to either limitations in the capability of the machine or Insufficient heating at the particular parameters; less extreme values for certain parameters were chosen in these cases. The different levels for the four parameters that were varied, along with parameters for the all high and all low welds, are shown in Table 1. The burn-off (displacement in the direction of applied force) and the applied force were monitored during welding. Data was gathered at 10 Hz during the welding process. From this data the burn-off rate (rate of material expulsion) and the duration of the frictional phase (friction time) were determined. The burn-off rate was determined by fitting a straight line to the burn-off data during a period in which the process is operating at steady state (Fig. 3). The fit was conducted using the least squares method, and the gradient of this line of best fit was taken to be the burn-off rate. The friction time was determined to be the period from the first contact of specimens until the oscillation had stopped (Fig. 3). 2.3. Characterisation

• frequency of oscillation: in this study it is defined as the number of sinusoidal oscillations (of the reciprocating part) completed in a second • amplitude of oscillation: defined as the maximum displacement of the oscillating sample from its equilibrium position (equilibrium position is when the displacement between the oscillating and stationary sample is zero, i.e. samples are aligned) • friction pressure: the pressure that is applied, perpendicular to the weld interface, during the rubbing, or frictional, phase of the process. Pressure is calculated by using the nominal area of contact

A selection of welds were sectioned for metallographic examination by cutting parallel to the reciprocating direction (Fig. 1). These sections were then mounted in Bakelite, ground and polished. Polished specimens were subsequently electrolytically etched in 60% HNO3 for a general analysis of the microstructure, or electrolytically etched in KOH for a quantitative analysis of the weld line delta ferrite fraction. The delta ferrite fraction was determined by image analysis using Image J software [25]. Tensile tests were conducted on a selection of welds in accordance with BS EN

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Fig. 3. An example of weld data from a linear friction weld in 316L. The figure shows how the burn-off rate and the duration of the frictional phase were determined. (A) The two specimens move to come into contact. (B) Contact occurs – start of frictional phase. (C) Duration of the frictional phase. (D) Oscillation is stopped and therefore there is very limited further burn-off; end of frictional phase. (E) Process reaches a steady state and material is expelled at a fairly constant rate. Burn-off rate is determined from the data in this period.

10002-1:2001. The tensile specimens had a diameter, D, of 12.5 mm and a gauge length of 5D. A texture analysis was completed on one weld, using high energy synchrotron X-ray diffraction at the European Synchrotron Radiation Facility (ESRF) in Grenoble, France. The analysis procedure followed that of Romero et al. [26]. The sample was cut to a thickness of approximately 2 mm in the transverse direction (Fig. 1); this thickness was sufficient to allow diffraction images to be taken in transmission. A monochromatic beam of wavelength 0.142 A˚ (88keV) and 200 ␮m × 200 ␮m in size was used to record diffraction patterns at locations along a line going across the weld line at mid-width. A step size of 100 ␮m was applied, which resulted in overlapping diffracting gauge volumes. A high resolution Pixium 4700 area detector, placed 646 mm away from the sample, was used to collect complete Debye Scherrer diffraction rings. With the given energy/sample–detector distance it was possible to obtain a full representation of the austenite texture by using a single diffraction image. In order to fully represent texture an orientation distribution function (ODF) is needed, in which the coordinate system of the crystal and the sample are aligned through the use of three angles (ϕ1 , ˚, ϕ2 in Bunge notation) [27]. The main steps involved in obtaining an ODF, from the two dimensional diffraction images collected, were:

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1. The sample-to-detector distance, the centre of the diffraction image and the tilt of the detector were calibrated through the measurement of diffraction patterns of an aluminium powder of known lattice spacing. The calibration was carried out using the ESRF software package Fit2d [28]. 2. Each diffraction image was sliced around 5◦ angles to produce 72 different sections, using Fit2d [28]. In each of these sections the integrated intensities were calculated to produce a plot of diffraction angle, 2, against intensity. 10 austenite intensity peaks (corresponding to 10 different planes of the austenite phase) were available for analysis, over a 2 range between 2.5◦ and 13.7◦ (q-range 3.9–21 inverse angstroms). 3. The peak intensities from the diffraction spectra, in each 5◦ slice, were extracted using the Le Bail method [29], as implemented in the software package MAUD [30]. Incomplete pole figures, for each of the 10 planes analysed, were extracted from this data. The E-WIMV [31] algorithm in MAUD was then used to calculate an ODF from the incomplete pole figures, and complete pole figures were calculated from the ODF obtained.

3. Results and discussion 3.1. Macroscopic examination Visual examination of the welds showed most to have ample flash coming from the weld interface (as in Fig. 4a), which is a characteristic of welds of high structural integrity. In all cases the flash was bifurcated (the flash coming from the two weld halves splits and forms two separate collars, Fig. 4a–c), which has not been observed when LFW titanium alloys, where a single wing like structure is formed [3,21]. Welds that were representative of the different parameter sets were sectioned for detailed metallographic characterisation. Although most of the sectioned samples showed a complete bond line with no visible defects (as in Fig. 4a), there were a limited number of welds with defects caused by a lack of bonding. An incomplete bond was found when applying a low pressure (80 MPa) during welding, with bonding taking place at the centre of the weld interface, but not towards the edges (Fig. 4b). A lack of bonding across much of the weld line was found in the weld produced with all parameters set to the low level (Fig. 4c). In both of these cases the defects are thought to be because of the poor consolidation of welds, demonstrating that there is a low parameter threshold, or minimum energy level, for obtaining structurally integral linear friction welds in austenitic steel. A similar energy threshold was previously observed when LFW Ti-64 [3].

Fig. 4. Macro shots of welds produced with (a) high pressure (S17), (b) 80 MPa applied pressure, all other parameters at mid level (S20) and (c) all parameters set to the low level (S23).

Parent Bond line Parent Parent Parent Parent Parent Bond line Parent Bond line 73 26 73 71 70 72 73 29 70 11 46 44.5 46.5 47.5 45 48.5 45.5 40 49 14.5 601 598 594 603 610 610 601 593 598 512 286 303 310 309 296 284 292 296 277 300 40 45 25 48 16 40 60 20 57 4 160 160 160 160 160 160 240 80 200 60 All mid High frequency Low frequency High amplitude Low amplitude Low burn-off High pressure Low pressure All high All low S1 S2 S3 S4 S5 S6 S7 S8 S9 S10

40 45 25 40 40 40 40 40 45 25

2.5 2.5 2.5 3 1 2.5 2.5 2.5 2.5 1

3 3 3 3 3 1 3 3 2 2

0.2% Proof stress (MPa) Max power input (MNmms-1) Burn-off distance (mm) Amplitude (±mm) Pressure (MPa)

Frequency (Hz)

3.2. Mechanical testing A number of welds were tensile tested, and in most cases cross weld tensile test failure occurred within the parent material region of the gauge length. This indicates that the weld region had a higher ultimate tensile strength than the parent material. This was the case in all welds apart from the poorly consolidated welds (i.e. applied pressure of ≤80 MPa and all low parameters) and a weld made at a high frequency of oscillation. A summary of the tensile test results is shown in Table 2. 3.3. Microscopic examination

No.

Table 2 Summary of tensile test results of 316L linear friction welds. A selection of 10 of the 66 welds were tested.

Ultimate tensile strength (MPa)

Elongation (%)

Failure location

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Reduction in area (%)

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The general microstructure of six 316L linear friction welds were examined. Fine grains were seen close to the weld line in most of the welds (see Fig. 5 for a typical microstructure). The area of grain refinement was quite large, spanning between 500 and 750 ␮m in the different welds. A considerable amount of ␦-ferrite (in the form of black stringers in Fig. 5) could also be seen close to the weld line in some welds (effects of welding parameters on ␦-ferrite formation is discussed further in Section 3.5). Further away from the weld line there was a region spanning between 180 and 450 ␮m (on either side of the grain refined region), in the different welds, where the microstructure had been deformed slightly but a large amount of grain refinement had not taken place. The areas of grain refinement and deformation are, collectively, commonly termed the thermo-mechanically affected zone (TMAZ) [2]. Beyond the TMAZ the parent material microstructure appeared to be sustained. Although this general description of the microstructure was accurate for most of the welds, one of the welds studied had a very different microstructure. The weld zone of S20 (see Table 3 for welding parameters) had grain sizes that were much bigger than in the other welds and there was no readily visible change in grain size going across the weld line and into the parent material. The reasons why this microstructure exists in S20 and why the zone sizes change with varying parameters is still an area of ongoing research. It is likely that the microstructures are produced as a result of a complex interaction between heat input, heat losses (due to conduction, convection and flash expulsion) and applied force. 3.4. Texture in a high pressure weld Fig. 6 shows the texture variation induced by the LFW process for the high pressure weld S17 (see Table 3 for welding parameters). Measurements from the weld line showed a strong {1 1 1} 1 1 2  type texture at the weld line (Fig. 6a), in which grains were predominantly orientated in the {1 1 1}1 2¯ 1 and {1 1 1}1¯ 1¯ 2 orientations (rotation of about 60◦ around the axial direction (ϕ1 ); Fig. 6d). There were slight deviations from these predominant orientations with small rotations around the transverse and axial directions (ϕ1 ), as shown in Fig. 6d. The same type of texture, but of reducing strength, was seen until just beyond the boundary between the grain refined and deformed regions (Fig. 5). Beyond this boundary the texture was still predominantly of the {1 1 1} 1 1 2  type, but there seemed to be a greater amount of rotation around the axial direction, than at the weld line (Fig. 6b). A similar amount of rotation occurred around the transverse direction. This type of texture continued to be present, but gradually reduced in strength, until a close to random texture was seen at a distance of about 3.3 mm from the weld line (Fig. 6c). As described earlier Chowdhury et al. [24] found a {1 1 1} fibre texture as a minor component in heavily cold rolled 316L, but to have a very strong {1 1 1} 1 1 2  component, without any other significant texture components, is quite unexpected. However, the {1 1 1} 1 1 2  type texture is a typical shear texture in face centre

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Fig. 5. Typical microstructure of a 316L linear friction weld (S17: high pressure).

cubic materials and has been reported as a prominent component in surface layers of rolled aluminium [32,33]. Mechanical twinning is thought to be a major factor in the formation of this type of texture. It has been observed during the cold rolling of 316L that lamellae of deformation twins cluster in bundles consisting of the twin lamellae and a thin lamellae of matrix material [34]. These bundles are orientated parallel to a {1 1 1} plane and slip predominantly continues on {1 1 1} planes parallel to the deformation twins [35,36]. This is thought to be because of the difficulty of slip on planes intersecting these tightly packed twin bundles (i.e. latent hardening occurs on planes other than those parallel to the deformation twins) [35]. The predominant slip on just one plane is thought to align the {1 1 1} plane towards the rolling plane during rolling and similarly towards the weld plane, giving {1 1 1} type textures, during LFW. 3.5. Effects of welding parameters on ı-ferrite formation As mentioned previously, in some welds a significant increase in ␦-ferrite could be seen close to the weld line. However, on

analysis of a number of different welds (S11 to S24) large differences in the fraction of ␦-ferrite close to the weld line were observed. The parent material contains about 1–2% ␦-ferrite, and this was maintained in the welds made with all parameters set to a low level (S23 and S24) and in the low frequency welds (S12 and S13, Fig. 7a). In contrast, the welds made at low pressures contained about 8% ␦-ferrite (≤80 MPa, S20 and S21, Fig. 7b). The fact that the microstructure and ␦-ferrite fraction of linear friction welds in 316L can be controlled by controlling welding parameters may make the LFW process advantageous for certain applications (e.g. applications involving long term high temperature exposure, where there is a susceptibility of ␦-ferrite decomposition to ␴ phase). As ␦-ferrite is a high temperature phase a lesser amount of it should be present in welds where the welding parameters produced cooler welds. An important factor affecting ␦-ferrite formation appears to be the burn-off rate (rate of material expulsion into the flash). Of the 66 welds produced in this study, the low frequency welds (S12 and S13, which had low ␦-ferrite fractions close

Table 3 Summary of data from the analysis of 316L linear friction welds. Data for 14 welds, of the 66 produced, are presented in this table. PQ indicates poor quality welds with microstructural defects. No. S11 S12 S13 S14 S15 S16 S17 S18 S19 S20 S21 S22 S23 S24

All mid Low frequency Low frequency High amplitude Low amplitude Low burn-off High pressure

Low pressure All high All low All low

Pressure (MPa)

Frequency (Hz)

Amplitude (mm)

Burn-off distance (mm)

Max power input (MNmms-1)

Burn-off rate (mm/s)

Friction time (s)

Weld line delta ferrite fraction (%)

160 160 160 160 160 160 240 120 100 80 60 200 60 60

40 25 25 40 40 40 40 40 40 40 40 45 25 25

2.5 2.5 2.5 3 1 2.5 2.5 2.5 2.5 2.5 2.5 2.5 1 1

3 3 3 3 3 1 3 3 3 3 3 2 2 2

40 25 25 48 16 40 60 30 25 20 15 57 4 4

2.3 4.1 4.3 2.5 3.2 2.7 3.2 1.9 1.4 0.9 0.6 2.8 0.4 0.7

2.3 2.3 2.3 2.4 2.8 1.5 2 2.9 3.5 4.8 6.8 1.7 15.2 12.1

7 2 2 6 1 7 6 7 6 8 8 4 2 1

PQ PQ PQ PQ

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Fig. 6. (a) Pole figures from the weld line of S17 (a weld produced at high pressure), showing a strong {1 1 1} 1 1 2  type texture. (b) Pole figures from the region at the boundary between the grain refined and deformed zones (∼300 ␮m from weld line, Fig. 5). (c) Pole figures at 3.3 mm from the weld line showing an almost random texture. (d) Ideal poles for the {1 1 1} 1 1 2  type texture and how (a) differs slightly from these. (e) Orientation of the pole figures.

Fig. 7. (a) Weld produced at low frequency (S12) with a similar amount of delta ferrite, at the weld line, as in the parent material. (b) Weld produced at low pressure (S21) with a significant increase in delta ferrite fraction at the weld line, relative to the parent material.

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force, F (Eq. (1), Appendix A). Pm = F × vm = F × 2f˛

Fig. 8. Variation of weld line delta ferrite fraction with burn-off rate.

to the weld line) were produced with some of the highest burn-off rates. In high burn-off rate welds hot material from the interface would be expelled out of the weld more quickly and this should mean there is less time for material close to the interface to heat up. A consequence of this is that the weld is exposed to a high temperature for only a short time, and therefore there is little chance of ␦-ferrite formation. A total of 14 welds were studied for weld line ␦-ferrite fractions and Fig. 8 shows that there is a linear relationship, in the studied range, between ␦-ferrite fraction and burn-off rate. Fig. 8 confirms that the short time at high temperature experienced by high burn-off rate welds has a significant effect on the ␦-ferrite fractions present in the weld region. The welds that did not fit the pattern were those produced with low level welding parameters (circular data points), which were also of poor quality, and those produced with a low oscillation amplitude (triangular data point).

3.5.1. Discussion of welds that did not obey a linear relationship between burn off rate and ı-ferrite fraction There were three welds that did not fit into a linear relationship between burn-off rate and ␦-ferrite fraction at the weld line, and these welds will be discussed in this section. Another possible mechanism for low ␦-ferrite formation would be a low welding temperature, and this appears to be the case for the welds produced with all parameters set to the low level (S23 and S24). A comparative assessment of the maximum power input, Pm , can be obtained by multiplying the maximum sliding velocity, vm , by the applied

(1)

where f is the frequency and ˛ the amplitude. Through this calculation it can be seen that the power inputs in S23 and S24 were significantly lower than in the other welds (Table 3). This suggests that a particularly low welding temperature was achieved in these welds. Furthermore, the low heat input and welding temperatures were not sufficient to produce welds of high structural integrity, as seen from the very poor mechanical properties in Table 2 (S23 and S24 were made at the same parameters as S10). The remaining anomalous weld in Fig. 8 is S15 (produced at the same welding parameters as S5, Table 2), which was produced with a low oscillation amplitude. A moderately large pressure was applied to produce this weld, together with a relatively low power input. Hence two factors resulting in a cool weld (high material ejection rate and low power input) were applied simultaneously resulting in an exceptionally low fraction of ␦-ferrite. Although Eq. (1) is useful for explaining the behaviour of the anomalous welds in Fig. 8, it is difficult to use it to determine trends, between ␦-ferrite and power input, across the whole of the parameter space. This is largely because the power input and burn-off rate are interdependent but tend to give opposite effects on the final ␦ferrite fraction. For example, if the welding pressure is increased the power input increases encouraging higher temperatures. However, the burn-off rate also increases with welding pressure (as discussed further in Section 3.7) encouraging greater heat ejection and lower temperatures. It is therefore hard to predict the results from these competing influences without using a computer program to model the process, which, although would be very useful, is beyond the scope of this work. 3.6. Sigma phase As mentioned earlier, ␴ phase formation has been observed in some regions of friction stir welded 304 austenitic stainless steel [17], with its formation being explained by an accelerated decomposition of the ␦-ferrite phase into the ␴ and ␥ phases during cooling of the weld. For S17 (a high pressure weld), however, no ␴ phase was detected in the diffraction traces from the synchrotron X-ray experiment (Fig. 9). A number of welds (S13, S17, S19) were also tested for sigma phase by using optical microscopy, in accordance with test method A of ASTM A923-08, with no sigma phase detected in any of the welds. These welds were chosen because they covered a range of burn-off rates (from 1.4 to 4.3 mm/s; see

Fig. 9. Synchrotron X-ray diffraction pattern taken at the weld line of S17, showing detectable amounts of only austenite and delta ferrite phases. No sigma phase was detected in the diffraction pattern.

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Fig. 10. Effects of pressure, frequency and amplitude on burn-off rate. Data for 16 different parameter combinations and 48 welds are presented in this figure.

Table 3) and weld line ␦-ferrite fractions (from 2% to 6%), whilst also exhibiting no microstructural defects. 3.7. Weld data examination The variation of burn-off rate with different process parameters is explored in more detail in this section, as the burn-off rate has been demonstrated to have a significant impact on the microstructure of the weld. It was seen that burn-off rates increased almost linearly with welding pressure (Fig. 10a). The relationship with frequency and amplitude, however, was very different, and there seemed to be a peak in burn-off rates at mid range frequencies and amplitudes (Fig. 10b and c). This relationship is similar to that seen with different rotation speeds during rotary friction welding [37,38], and has been seen to exist in a number of different materials, including carbon steel [37]. The relationship between the pressure and the burn-off rate is not unexpected since applying high weld pressures will expel material more easily and at an increased rate. However, there does appear to be a limit as to how high burn-off rates can increase with increasing pressures. This is evident in the slight curvature of the burn-off rate/pressure curve in Fig. 10a, but is much clearer from an analysis of friction times (Fig. 11). The friction time is closely related

to the burn-off rate as a set burn-off distance was used to trigger amplitude decay. The flattening off of the curve at high welding pressure shows that there is a limit to how fast material can be expelled from the joint with increasing pressure. The effects of frequency and amplitude on burn-off rate may be a result of the differences in contact conditions, at the interface, with different welding conditions. At low frequencies and amplitudes the burn-off rate would be expected to be low because the low heat input would create a small amount of plasticised material and therefore produce a low burn-off rate when it is removed [37]. At very high frequencies and amplitudes a layer of very soft material is thought to form at the interface because of the high heat input [39]. This material essentially lubricates the joint and is very hard to remove, which causes the low burn-off rate. Furthermore, the slow moving layer of hot material allows enough time at high temperature to produce a large amount of ␦-ferrite close to the weld line, as is evident from Fig. 8. In contrast, large layers of plasticised material can be removed from the joint almost as soon as they are formed at mid-range frequencies and amplitudes. This gives high burn-off rates as well as very short times at high temperatures, which limit ␦-ferrite formation (Fig. 8). Frequency and amplitude will both act in a similar way in increasing heat input, and therefore have similar impacts on burnoff rate, as both parameters affect the sliding velocity of the oscillating part (the sliding velocity along with the applied pressure being the key factors in producing heat input). Fig. 12 highlights this point by again showing a peaked curve when plotting burn-off rate against maximum sliding velocity. However, there appears to be a slight dip in the curve at mid range sliding velocities (at about 400 mm/s). This indicates a slightly more complex relationship, and suggests certain combinations of amplitudes and frequencies are more effective at removing material than others, even if the amplitudes and frequencies combine to give the same sliding velocity.

3.8. Effects of applied pressure on the effective welding temperature

Fig. 11. Effect of welding pressure on friction time. Data for 8 different parameter combinations and 24 welds are presented in this figure.

As mentioned previously, studies on the effects of welding parameters, especially applied pressure, on welding temperature have produced some contradictory results [19–21]. A number of possible explanations have been put forward which may account

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change in solely the welding time without the need to change other more significant parameters, e.g. applied pressure, frequency, etc., that can influence the power input and output. When comparing S16 and S11 (Table 3) it can be seen that both welds had very similar levels of ␦-ferrite at the weld line even though S11 had a higher burn-off distance, and therefore a longer welding time (or friction time as heat is only generated in the frictional phase). This shows that the welding time has no significant effect on the peak welding temperature and weld line ␦-ferrite fraction. Since most welds reach a steady state condition during LFW, where the power input is balanced by the power output (i.e. because of flash expulsion, conduction and convection), this observation seems sensible. 4. Conclusions Fig. 12. Effect of maximum sliding velocity (a function of the frequency and amplitude) on burn-off rate. Data for 8 different parameter combinations and 24 welds are presented in this figure.

for these discrepancies, however the results from this study may help to identify another possible explanation. The findings from the present work suggest that on application of parameters that give low power inputs the applied pressure contributes largely towards producing the frictional heat input. In these circumstances an increase in welding pressure will produce a higher welding temperature. Conversely, when using welding parameters that give high power inputs the pressure largely helps to expel hot material, and increase the burn-off rate. Consequently, in this circumstance, increasing welding pressure causes an increasing amount of material and heat ejection and therefore a cooler weld is produced. Hence, the pressure under different conditions can influence the temperature of the weld in very different ways, which may have led to the discrepancies described for linear friction welded titanium alloys. Although the welding temperature has been suggested to reduce with increasing applied pressures, it is important to stress that a time at temperature effect, rather than an actual lowering of temperature, may cause the microstructures seen at high pressures. In the high pressure welds a narrow band of hot material is expected to form as there is little time to transfer heat from the material at the interface to the surrounding material. This is because the high burn-off rate in high pressure welds largely restricts the time the material is exposed to high temperature. Much of this very hot material will be ejected from the weld (into the flash) at the end of the process because of the forge pressure, which is applied as the material is cooling. This would leave a weld that appears to have been produced at a low temperature because much of the very hot material from the weld line has been expelled. There is also the possibility that a high welding pressure will allow material to be ejected into the flash at relatively low temperatures [19,20]. However, this would not explain why 316L welds made at the same pressure, but with different amplitudes and frequencies, produced very different ␦-ferrite fractions (ranging between 7% and 1%, Table 3) at the weld line. If different applied pressures were causing material to be expelled at different temperatures a similar level of ␦-ferrite would be expected in welds made at the same pressure. This is because the sliding velocity is not likely to have a large effect on the temperature at which material is expelled. This reasoning leads to the suggestion that a time at temperature effect is the dominant factor in producing variations in microstructures with different applied pressures. The results of the present study also suggest that a change in burn-off distance has no effect on the maximum welding temperature. The variation of burn-off distance effectively allows a

Sound linear friction welds, with cross weld tensile test failure in the parent material, could be produced in AISI 316L. The process, with regard to welding parameters, was demonstrated to be quite robust and fairly large changes in welding parameters did not, in most cases, affect the tensile strength, ductility or defect levels of the weld. A strong {1 1 1} 1 1 2  type texture was seen at the weld line of a 316L linear friction weld. This is a typical shear texture in face centre cubic materials. Large variations in the fraction of ␦-ferrite close to the weld line have been observed in welds produced by using different welding parameters. The control of ␦-ferrite fraction by just varying welding parameters may make the LFW process advantageous for certain applications. A number of general conclusions were drawn from studying the effects of welding parameters on the welding process and on the final ␦-ferrite fraction. • The burn-off rate has a very large effect on the final ␦-ferrite fraction, with high burn-off rates producing low levels of ␦-ferrite. This is because material close to the weld line experiences a short time at high temperature, in high burn-off rate welds, as hot material at the interface is ejected into the flash more quickly. • The highest burn-off rates occur when applying a high welding pressure in combination with mid-range amplitudes and frequencies (about 30 Hz and 1.5 mm amplitude in 316L). • High burn-off rates are not thought to result in a decrease in the welding temperature, but enable the creation of a narrow band of hot material that is more readily ejected into the flash during the forge phase. Therefore material that is exposed to relatively low temperatures is present at the weld line on completion of the weld. • A sole change in welding time, achieved through variation of burn-off distance, does not appear to have a significant effect on producing different microstructures with different welding parameters. Although these general conclusions were drawn from studying linear friction welded 316L, they should be applicable, possibly with some small to moderate variations, when welding other engineering alloys, such as typical aeroengine materials. Acknowledgements This work was partially funded by the industrial members of TWI Ltd. as part of the Core Research Program (CRP), with financial support also provided by the EPSRC and TWI Ltd. The authors would therefore like to thank these bodies. The authors would also like to acknowledge the contribution of many staff within the friction and materials departments of TWI, and thank the ESRF for access to the ID15B beamline.

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Appendix A. Determination of the relative power input P = Ff × v = Ff × ˛ω cos(ωt) = Ff × ˛2f cos(2ft)

(2)

where P is the power input, Ff is the maximum force parallel to the oscillation (i.e. in the reciprocating direction), and ω is the angular frequency, 2f, where f is the frequency. The maximum value that the cos(ωt) term in Eq. (2) can achieve is 1, which results in an equation for the maximum power input, Pm (Eq. (3)). Pm = Ff × ˛ω = Ff × vm

(3)

The force in the reciprocating direction can be related to the applied force, F, through a friction coefficient,  (Eq. (4)) [40]. Ff = F

(4)

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