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The application of neutron diffraction to a study of phases in type 316 stainless steel weld metals
165
Journal of Nuclear Materials 118 (1983) 165-178 North-Holland Publishing Company
THE APPLICATION OF NEUTRON DIFFRACTION TO A STUDY OF PHASES IN TYPE 316 STAINLESS STEEL WELD METALS G.F.
SLATT’ERY
UKAEA,
Risley Nuclear Power Development
Laboratories,
Risiey, Warrington
WA3 6AT, UK
and C.G.
WINDSOR
UKAEA, Materials Physics Division, Atomic Energy Research Establishment, Received
21 April
1983; accepted
Hanvell, Oxon, OX1 I ORA, UK
6 June 1983
Neutron diffraction techniques have been utilised to study the phases in type 316 austenitic stainless steel weld metal, both in the as-welded condition and after stress-relieving and ageing heat-treatments. The amounts of the principal crystallographic phases present in bulk specimens have been measured. Two compositions of weld metal were selected to provide a “low” (6%) and “high” (16%) initial ferrite level and the subsequent volume fractions of transformation products were measured after heat-treatment. Some retained ferrite was observed in all the heat-treated specimens, ranging from 4% for specimens of both initial ferrite levels treated at 625’C for 1000 h, to around 1X for the specimens treated at 85O’C for 6 h. The high initial ferrite specimen produced 0.9% of sigma phase after the 850°C treatment and 0.2% sigma after the 625°C treatment. The low initial ferrite specimen produced 1.5% M,,C, carbide after both heat-treatments. The results compare well with previous findings on similar samples of weld metal using optical and electron microscopy.
1. Introduction Welds in type 316 austenitic stainless steel are of considerable interest because of their usage in both conventional and nuclear power stations. As a result, a high standard of structural integrity is expected from the weldment. The composition of the weld is formulated so as to provide between about three and eight per cent of delta-ferrite in the microstructure to avoid hot tearing problems during weld solidification [l-4]. The amount of S-ferrite present is essentially dependent on the chemical composition of the weld deposit and various combinations of chromium, nickel and molybdenum are used in the weld metal to provide the optimum amount of S-ferrite. Typical compositions of the weld metal used in service are 19Cr12Ni3Mo and the leaner 17Cr8Ni2Mo [5]. In practice, the weld metal is deposited as a succession of overlaid weld beads. The macrostructure of such multipass weld layers in influenced by the absence of any phase transformation in the weld on cooling, with the result that the structure consists of long, columnar
0022-3 115/83/0000-0000/$03.00
0 1983 North-Holland
grains which nucleate from favourably oriented grains in the heat-affected zone and grow epitaxially from bead to bead in the weld [5]. The columnar grains grow along lines of heat dissipation such that a strong elastically anisotropic character is developed in the weld metal. This feature of marked anisotropy renders the ultrasonic inspection of such welds difficult owing to directional variability of the velocity and attenuation of the elastic waves [6,7]. Each grain contains a network of S-ferrite distributed throughout the austenite matrix such that the fine-scale structures consists of a series of low-angle boundaries or dislocation networks linking the S-ferrite into a sub-grain structure with a diameter of less than 10 pm [8]. Large components of weld metal need to be stressrelieved for service usage to reduce or even eliminate the residual stresses introduced during the welding cycle. During such heat-treatment, the S-ferrite phase in the weld microstructure transforms to produce M&, carbide and intermetallic phases such as chi and sigma [8-111. Since the S-ferrite is enriched in chromium and molybdenum due to microsegregational effects [ 12-141
166
G.F. Slattery C.G. Windsor / Neutron drffaction study of starnless sreel
then these transformations occur much more rapidly in the ferrite phase or at the ferrite-austenite interface than in the austenite matrix. Consequently, these reactions take place more readily in a duplex weld structure than in the equivalent wholly austenitic parent steel. Since the production of these phases on heat-treatment affects both the chemical and mechanical properties of the weld metal, it is important to understand how these properties may be modified by changes in composition and microstructure of the weld metal. Thus changes in composition of the weld metal can affect not only the amount of S-ferrite present in the original weld but also the amount of intermetallic phases formed as a result of the transformation of the S-ferrite during subsequent heat-treatment [5]. A study was initiated to investigate the influence of different levels of a-ferrite on the phase transformations produced as a result of both a stress relieving treatment at 85O’C and also an ageing treatment at 625°C to simulate an in-service usage. Since previous work has been carried out on the structure of type 316 steel both in the wrought and welded conditions using primarily electron-microscopical techniques [ 12-201, a detailed study was first made on these stress-relieved and aged welds using techniques of microanalysis with the electron microscope together with selective etching methods and optical microscopy [21]. This work is being reported separately [22]. The second stage of the work reported here, utilises the alternative technique of neutron-beam diffraction which is a complementary technique to more conventional X-ray diffraction. The study assesses the feasibility of using neutron diffraction on bulk samples of significant thickness to provide the bulk volume fractions of the phases in the welds in order to correlate the results with the previous electron/optical microscopic studies on the identical samples of weld material. Neutron diffraction has the great advantage that the beam is able to penetrate metals since, having no electric charge or electric field, it interacts only weakly with matter [23]. Neutron beams can therefore be used to study bulk effects for example in steels, the penetration depth is approximately 8 mm. Also the surface condition of the sample is unimportant so that no special preparation techniques are needed. In contrast, electron beams interact much more easily and so penetrate essentially only comparatively small distances. In the present study, neutron diffraction was used (1) to measure the volume fraction of S-ferrite originally present in the weld samples together with the amount of untransformed S-ferrite after heat treatment, (2) to obtain basic crystallographic information on the face-centred cubic y-austenite and body-centred cubic S-ferrite phases and
(3) to identify and quantify the amount of the phases produced as a result of the transformation of d-ferrite during heat treatment.
2. Experimental procedure 2.1. Materials and heat-treatment Two commercially available type 316 welding electrodes to B.S. 2926: 1970 were chosen to give approximately 5 and 15% of b-ferrite in the weld metal. Their compositions and actual ferrite levels are given in table 1 which shows that the welds had nominal compositions lKr9Ni2Mo and 20Cr lONi3Mo. The increased levels of the ferrite-forming chromium and molybdenum in the second composition resulted in the higher level of d-ferrite content. The weld deposits were laid down by manual metallic arc-welding to give weld pads 150 x 25mm in size typical welding conditions being: current, 125 A; voltage, 20 V; electrode size, 4 mm diameter; heat input, 1 kJ; with maximum interpass temperature, 150°C. Specimens were examined in the neutron beam diffractometer in the as-welded condition and also after the following heat treatments, namely (1) stress-relieved at 850°C for 6 h and (2) aged at 625°C for 1000 h. Microscopical examination showed that as-welded the “low” ferrite steel (6%) possesses a discontinuous network of ferrite with a vermicular type morphology [24,25] as the minor phase distributed within an austenite matrix. In the “high” ferrite weld (16%), the a-ferrite formed a more continuous network characteristic of the morphology designated as “lacy” in type [24]. Also present in the general microstructures were small spheroidal slag particles approximately 0.5 to 3 pm in
Table 1 Composition (wtW) and ferrite content of welds Element
“High” ferrite weld metal
“Low” ferrite weld metal
Carbon Manganese Silicon Chromium Nickel Molybdenum Phosphorus Sulphur Ferrite volume fraction (%)
0.047 0.96 0.41 20.5 10.1 3.13 0.02 0.01
0.067 2.18 0.35 18.0 8.9 1.85 0.02 0.01
15.6
5.9
167
G.F. Slattery, C.G. Windsor / Neutron diffraction study of stainless steel
Fig. 1. Scanning electron micrograph of a heat treated d-ferrite network of the duplex weld metal.
weld sample showing
precipitation
of sigma and chi intermetallic COUNTER
phases in the
SHIELO I
size. Fig. 1 is a scanning electron microscope image of a heat-treated weld sample showing precipitation of sigma and chi intermetallic phases in the S-ferrite network of the duplex weld metal. 2.2. Neutron
MONOCHROMATOR
diffraction
The experimental technique is very similar to X-ray diffraction with the substitution of the nuclear reactor for the X-ray source. Neutrons of a particular wavelength X are selected from the Maxwellian distribution of neutron wavelengths present in the core by Bragg reflection from a single crystal. The beam direction is defined by means of a slit collimator, and the beam size adjusted so that it matches the sample size. The scattered neutrons are detected by one or more gas counters. The scattered neutron intensity is measured as a function of
-_____ Fig. 2. The layout reactor, Harwell.
~of the CURRAN
diffractometer
at DID0
BEAM
STOP
v
G.F. Slatrery, C.G. Windsor / Neutron diffraction study of stainless steel
168
28 = 90”. The instrumental peak width was fitted to the formula W= A sin 8/AB with A = 806”/nm and 3 =
the scattering angle “A” to plot out the diffraction pattern. Fig. 2. shows the layout of the CURRAN diffractometer in the DID0 reactor at Harwell used in these measurements 1261.The monochromator used was hot pressed “squashed” germanium at a take-off angle 26, = 47.5’. Using the (331) germanium reflection gave a wavelength of (1.369 f 0.002) A as determined from a nickel powder calibration scan. This wavelength is convenient in that the second-order scattering (h/2) from the mon~hromator is negligible, while the range of d spacings down to about 0.1 nm covered by the scan covers several austenitic and ferritic peaks. The counter bank contains 5 distinct counters with approximately 10” separation. The exact separation and the nominal zero of the scattering angle is found from the nickel powder scan. The spectrometer was used with a collimation of 0.5” before the sample and about lo geometric collimation after the sample. This gave peak widths of order 0.4’ giving a resoiution of Ad/d = 0.008 near 28 = 4S”, rising to a width of 1.5Oor Ad/d = 0.014 near
RIS7
2.41. Relative intensity between the counters was calibrated by means of a previous calibration run using a vanadium incoherent scattering sample. 2.3. Sample preparation The samples for diffraction examination were sawn from the weld pad. They were approximately 27 mm high, 15 mm wide and 2 mm thick. They were mounted on the spectrometer with the base of the sample pressed into a piece of plasticene, covered in cadmium foil to act as a neutron absorber. This was in turn mounted on a base which rotated at about 4 revolutions per minute. This rotation was made in order to reduce preferred orientation effects in the sample. Since these were not eliminated by rotation in only one axis, the results will still contain an inaccuracy from residual preferred orientation effects. The sample was centred in the neutron beam by taking photographs using a scintillator plate holder.
HIGH
6H850
6O.OC
0.00
5.00
1
I
15.00
25.00
35.00
45.00 2-THETA
Fig.
3.
I
,
,
55.00
65.00
75.00
fOEGREEt
An example of the raw diffraction pattern on a scale to include the matrix peaks.
L
85.00
1
95.00
G. F. Slattety,
C.G. Windsor / Neutron diffraction study of stainless steel
r: c I
t
RIS
11
I
i
IGH
6HE.O
fig. 4. The raw neutron diffraction patterns on an expanded scale. The zeros of the graph are suppressed and no normalisation has been performed. Most of the runs lasted about 20 h and corresponded to a scan of from 7” to 102” in 0.2’ steps. The counting time at each point was varied to correspond to fixed
counts on a fission chamber monitor counter placed in the incident beam. By this means, variations in the incident flux are cancelled out. A separate run was
G.F. Slattery, C.G. Windsor / Neutron diffvuction study
170
made without a sample in order to determine the instrumental background. This turned out to be small compared to the disorder scattering from the several components of these steels together with incoherent scattering from the sample at ali angles. A specimen of nickel powder was mounted and scanned in the same way as
Table 2 Interplanar Probable assignment
d spacings
of observed
reflections
the other specimens in order to determine zero angle and the wavelength. 2.4. Data processing
Data from each run were punched on paper tape which was analysed on a DEC PDP 1 l/40 minicom-
1000 h at 625°C
6 h at 850°C
High
Low
High
Low
1.0183(g)
1.0844(5)
1.9394(5) 1.0857(2)
1.O 124(30) 1.0379(5) 1.0841(2)
1.0379(5) 1.0851(2)
I .0380(4) 1.0841(27)
a 211 (1, 720 y 220
1.1742(6)
l.i780(3)
1.1712(7)
1.1728(3)
I. 1728(7)
1.2770(2)
1.273415)
1.2717(4)
1.2727(2)
1.2717(2)
cr c
200 640
1.4380(4)
1.4380(2)
1.4348(28) 1.4812(42)
1.4364(7)
1.4367(12) 1.4842( 17)
c d (I c7
622 501 322 312
Y y 0
200 222 331
cy 220 y 222 y 311
1.03&l(8)
1.5627( 17) 1.6111(30)
0 a y G c d
1.4927(U) 1.5505(6)
I .6066(45) 1.6388
1.6485(6) 1.6742(3)
I .7980(2)
I .8007(2)
1.7995(2)
320
rJ
310
c 0 c
220 101 200
1.7980(2)
1.7992(2) 1.8455(16) I .8931(4)
1.8893(17)
2.0340(5) 2.0766(3) 2.1267( 180)
2.0364(5) 2.0787(3)
2.0300( 10) 2.0761(3)
I .88 I7(59) 1.9373(4) 1.9827(5) 2.0315(5) 2.0776(3) 2.1327(27)
1.9377(4) 1.9841(5) 2.0330( 10) 2.0756(3)
2.0364(5) 2.0771(3) 2.1404(11)
2.1753(29) 2.3423(67)
0
1.0094(2) 1.0388( 1) I .0848(2) 1.1199(5) 1.1481(8) 1.1764(6) 1.2392(33) I .2724(2) 1.3315(6) 1.3837(15) 1.4360(9)
1.6060(60)
1.8813(13) 411 212 110 111 410 422 002
High
1.7659(22) 1.7988(4)
C440 0
the counter
(A) a
As welded Low
of starnless steel
2.3423(67)
2.3464(134) 2.6806(35)
2.9528(2 16)
2.9474(32)
2.9495(86) 3.7165(87) 5.28
1l(600)
5.7794(127)
2.2902( 19) 2.3551(20) 2.4447(7) 2.6648( 123)
2.3332(66) 2.3330(66)
2.2966( 13) 2.3558(20)
2.9593(43)
2.9420(2 15) 3.7287( 192) 4.1396(43) 5.3059(36) 5.8006(980)
2.6771(9) 2.7701(29) 2.9528(43) 3.6771(51) 4.0635(21) 5.2357( 173) 5.8609(44)
a The numbers in brackets refer to the statistical variance of the least squares fit to the d spacing. Absolute of 0.02% because
of systematic
errors
in the wavelength
calibration.
errors are larger by order
171
G.F. Slattety, C.G. Windsor / Neutron diffraction study of stainless steel
From the fitted peak scattering angles, the d spacing of each peak was calculated from the Bragg equation. The d spacings of all the located peaks are given in table 2 together with an estimate of the statistical error. In nearly all cases, the width of the reflection was slightly greater than the instrumental width given by the equation
puter. The first program PATN4 [27] takes the data from all the six counters and combines them together into a single file corrected for instrumental background zero angle shifts, counter efficiencies, and counter separations. Fig. 3 shows the peaks from the austenitic matrix on a scale to fit the graph. Fig. 4 shows the results on a greatly expanded scale to emphasise the minority phase peaks. A second program ROBOT locates the peaks automatically and finds their appropriate position and intensity. The third program CRAWL takes this file and refines the peak intensities, positions and widths over and above the resolution. The program also gives an estimate of the statistical errors in these variables. FERRITE 200
T
5 the
the
extra widths are shown above widths and expressed in the form
cam
8628
2~
2dtane
I
OD
//
628
(3)
.
A slope on the figure corresponds to strain broadening. It is such that both the austenite and ferrite peaks have a crystallite size of around 60 nm but that the ferrite peaks appear to have some additional strain broadening.
1
AS WELDED LOW FERRITE
(2)
’
2x
where A 28 is the extra width of the angular scan. In this figure, a flat line corresponds to a crystallite size.
500
I
fig.
L=Q=
I
0
(1)
AQ -= 4s
I 211
__+__'---I--
t
8,,wn/2.
In
_---
0 006
COS
instrumental
PEAKS
I
i 110
D= d
4
3. Results 3.1. Analysis
--
0
f
1000 HOURS AT 625'C LOW FERRITE I
ce ,
/
.* -
of the austenite and ferrite lattice constants
Austenite has a face-centred cubic structure (fee). For cubic structures, the lattice constant can be derived for any peak of known Miller indices (hkl) from the equation a, = d( h2 + k2 + 12)“2.
0
0
I I
I
;K$RS FERRITE
AT .350°C 111 200 AUSTENITE 1
1
'I
f
6 HOURS AT 850°C HIGH FERRITE
,PEAKS 0.1 1 Sin
Table 3 Lattice constants and volume phases in weld metal
B/h
f 220 1
1 311 222 I 0.2
IOD
,2r=
Fig. 5. The line widths of the austenite and ferrite reflections, above the resolution value, expressed as a broadening factor Q (left hand scale) or as a crystallite size (right hand scale). The finite slope of the ferrite peak widths corresponds to some strain broadening.
Austenite “Low” ferrite “High” ferrite Ferrite “ Low” ferrite “High” ferrite Ferrite volume “ Low” ferrite “ High” ferrite
(4)
fraction
of austenite
and ferrite
As welded
1OOOh at 625°C
6h at 850°C
3.5972(5) 3.601 l(6)
3.5959(6) 3.5991(5)
3.5959(7) 3.5982(4)
2.8763(2) 2.8792(2) fraction (%) 5.9 15.6
2.8682(5) 2.8729(S)
2.8737( 12) 2.8778(51)
4.0 3.5
1.6 1.0
172
G.F. Slattery,C.G. Windsor / Neutron diffracrion study of stainless steel 3.2. Volume fractions of the ferrite phase Neutron scattering intensities depend on many factors but all of these can be readily evaluated so that absolute intensity measurements are straightforward to an accuracy of a few per cent (to achieve the highest accuracy careful corrections must be made for factors such as resolution, multiple scattering and extinction). In the present context, it is sufficient to obtain intensities relative to the austenitic matrix so that many factors cancel. The total intensity of a powder peak integrated over the angular scan is given by Bacon as follows [28]
AUSTENITE
1 NZ,,,[F,,,12X3L,exp(-2W)
(I=-
""7t
, LOW AS WELDED
V sin Bra sin 20
8772
HIGH FERRITE-t
FERRITE-! lOOOh 625%
FERRITE 6h 850°C
Fig. 6. The lattice constants of the austenite and ferrite phases in “high” and “low” ferrite specimens as welded and after heat treatments.
’
assuming N unit cells each of volume V and where Z,,, is the multiplicity of reflection h/cl. The Lorentz factor, LJ2rsin28, represents the fraction of the whole Debye-Scherrer cone intercepted by the counter. The last factor is the Debye-Waller factor representing thermal motions. Fhk, is the structure factor depending on the mean neutron scattering length b of atoms in positions R in the unit cell such that Fhk,=bexp(-iQR).
(6)
For a primitive structure cell, such as the austenite The five reflections 111, 200, 220, 311 and 222 corresponding to the austenite pattern were always clearly visible, and all gave consistent values of the lattice constant. The values given in table 3 and plotted in fig. 6 show the mean lattice constant, and its variance determined from these five reflections. In the case of the body-centred cubic (bee) ferrite peaks, the accuracy is somewhat less. Only three reflections 110,200, and 211 were covered by the range of the experiment. Of these the alpha (110) peak was only seen as a “wing” to the gamma (111) peak, and was also unresolved from the sigma (202) reflection. Also the alpha (220) was very close to the gamma (222) reflection. Thus in practice the alpha (110) peak could not be used in sigma-containing steels and only two or three alpha reflections were used in determining the ferrite lattice constant. The results in fig. 6 show a clear and significant change in lattice parameter with heat-treatment. The austenite lattice contracts on both the 850°C/6 h and 625O/lOOO h heat-treatments. The ferrite lattice shrinks appreciably, more especially for the 625°C/1000 h treatment. However the “high” ferrite (16%) lattice constants are consistently more expanded than the “low” ferrite (6%) lattice constants.
4~
with only one atom per unit and ferrite phases,
= b
and thus (7 sin 0 sin 28 Z,,,
NA3b2
exp( - 2W) = F
(7)
= ”
that is, a constant for any given set of reflections for a particular phase. The relative number of unit cells for two phases can therefore be found from the ratio of the constants C. For example, for the ratio of body-centred cubic to face-centred cubic phases, we have
N
Vbccb:,, Cbcc
N FCC
vfcc
bee -=---
b;_
Table 4 Neutron scattering
Cfcc .
cross-sections
Element
At. wt
b (IO-”
Fe Cr Ni MO Si C
55.8 52.0 58.7 95.9 28.06 12
0.95 0.352 I .03 0.69 0.42 0.665
cm)
173
G. F. Slattery, C.G. Windsor / Neutron diffrraction study of stainless steel Table 5 The matrix and secondary phases in type 3 16 steel weld metal: approximate compositions
Phase
Crystal structure
Austenite Delta-ferrite Ma sCg carbide Sigma Chi
face centred cubic body centred cubic cubic tetragonal cubic
AI A2 A2 D8b A12
Fe (wt%)
Cr
Ni
Ma
68 66 2s 58 57
18 23 65 28 23
9 5
2 4 5 8 13
The neutron scaitering iengths for the nuclei present in type 316 stainless steel are given in tabfe 4. The phases identified in this work are listed in table 5 and using the atomic concentrations measured for the austenite and ferrite phases given in this table, the mean scattering lengths can be computed as b kc = 7.81 X b fee =
IO-“mm
8.47 X lo- I2 mm
The Debye-Waher factors for pure iron and for type 316 stainless steel have been determined by previous diffraction studies to high values of sin e/X on the
LINAC pulsed neutron fusing the conventional Wailer factors) are B roc=0.47A2;
source. Appropriate values symbohsm for the Debye-
B,,,=0.29$.
The constant C was computed for each reflection and a mean taken. Table 6 shows experimental and calcuhrted line intensities for the austenite and ferrite reflections. It is seen from table 6 and from fig. 7 where the experimental and calculated line intensities are plotted on a logarithmic scale against angle, that the intensities are genera& in agreement but that the calculated and observed intensities are occasionally in
Table 6 Caktdated and observed line intensities for the austenite and ferrite f& refieetions Reflections (W
Materiai As welded
Austenite 111 200 220 311 222 Ferrite
110 200 211
ObLWNd
calculated observed caktiated observed calculated observed calculated observed calculated observed calculated observed calculated observed calculated
la00 h at 625OC
6 h at 85O’C
Low %
High 6
Low 6
High 6
Low 6
High s
53 46 48 26 20 26 24 39 4.5 12.2 43 32 9.8 8.4 II 23
48 51 71 29 20 29 25 44 3.6 1.4 122 96 35 25 33 69
49 34 25 19 14 19 18 29 4.5 8.9 22 16 3.5 4.1 6.1 11.5
29
84 76 86 43 30 44 42 65 6.2 20.2 21 15 3.8 3.9 5.0 IO.8
68 73 101 41 20 42 36 63 5.4 19.4 12.7 8.6 0.62 2.22 7.8 6.2
29 36 16 13 17 15 25 2.5 7.7 28 I2 3.5 3.1 8.3 8.7
G. F. Slattery, C.G. Windsor / Neutron diffraction study of stainless steel
174
cause of the systematic error resulting from preferred orientation. The results in table 3 show that the as-welded specimens have about 6 and 16% ferrite as the duplex constituent with austenite. With either original ferrite percentage, around 4% ferrite remains after the ageing treatment at 625°C for 1000 h and 1% remains after the stress-relieving treatment at 850°C for 6 h. 3.3. The identification of precipitate phases
1000
c
LOW
The phases identified by electron microscopical techniques in type 3 16 stainless steel weld metal [ 10,221 and listed in table 5 are all of known structure [ 151 so that it was possible to predict the approximate d spacings and relative peak intensities expected from each structure. The d spacings were calculated with the KDRREF programme [29] and the line intensities by the TAILS programme [30]. Fig. 8 shows the calculated scattering angles and intensities on a logarithmic scale and to the same scale as the experimental results in fig 7. By laying one over the other, it is straightforward to identify those phases having several peaks in the observed spectrum.
HOURS 625Y FERRITE
3
L
HIGH
FERRITE
3
4 3 2 1 10
20
30
LO
50
60
2 80
Fig. 7. Experimental and calculated line intensities for the reflections plotted on a logarithmic scale versus angle. The located peaks in the six specimens are indicated by solid lines. The dashed line and side mark indicate calculated intensities.
error by a factor of two or three. This is probably the result of the appreciable preferred orientation which was present in the sample. The average value over a number of reflections which is used in determining the volume fraction should be appreciably more accurate. Table 3, lower section, shows the volume fraction of ferrite compared to austenite evaluated using eq. (8). The number is strictly a number fraction N,,,/N,,, but since the atomic volumes are so nearly equal, it is equal to the number fraction within experimental error. The statistical error on the volume fraction is negligible, but an overall error of perhaps 20% must be allowed be-
-PHASE 2I
1
3
Y
I
I
1
PHASE
2 II
100
I
200
I
300
I
1
400
500
60"
29s
Fig. 8. The calculated intensities and scattering angles of the reflections from phases known to be present in the steels.
175
G. F. Slattery, C. G. Windsor / Neutron diffraction study of stainless steel
By this means, the sigma phase could be clearly identified in the “high” ferrite specimens heat-treated at both 85O’C and 625°C. It was also clear that the M,,C, carbide structure was present in the “low” ferrite specimens heat-treated at both temperatures. However no other phases could be identified in the observed patterns and none of the minority precipitate phases such as carbide and intermetallics could be identified in the as-welded specimens. The sigma-phase lattice parameters were deduced from the 002 and 410 reflections and the carbide phase parameter from the 440 reflection as given in table 7. The line intensities calculated with the TAILS program were then matched with the experimental line intensities in the cases given in the table. The
total intensity summed over the lines given were then used to calculate the volume fraction of the phase concerned. Once again the systematic error caused by preferred orientation leads to an error of the order of 20% in these volume fractions. In all cases, the volume fraction has been evaluated by multiplying the relative number of unit cells obtained from the TAILS intensities by the number of atoms per unit cell for each phase. In the “high” ferrite sample after the 850°C heat-treatment, the sigma phase volume fraction was 0.9% but in the sample with the 625’C treatment it was only 0.2%. In the “low” ferrite samples, the volume fraction of carbide after both heat-treatments was about 1.5%.
Table 7 A comparison of the observed intensities (upper figures) with calculated carbide and sigma phases in the heat-treated weld specimens
intensities
(lower figures)
for different
reflections
from the
Sigma phase high-ferrite specimen
Carbide (M,,C,) phase low-ferrite specimen Treatment:
1000 h at 625OC
6 h at 850°C
Treatment:
1000 hat 625°C
6 h at 850°C
Lattice a, (A) parameters
10.64
10.64
a0 (A) CO
8.8 4.60
8.8 4.593
Int.
Int. 171 430 90 74 87 92
(W
Int.
Int.
(hW
200
(122) (20) (121) (34) 70 15 71 128 46 16 71 103 52 59
(151) (36) 126 62
101
222 400 440 533 622 711,640
Ill 310
135 231
320 002
113 186 77 106
410 212 411 331
85 13 126 122 195 186 163 118 259 392 186 182
222 312 322 501 Volume fractions
N( c)/N(
y)
0.015
0.015
NAN
0.0022
589 1217 1910 1850 1907 1176 3282 3893 2155 1818 935 347 604 477 224 324 142 389 0.0091
G.F. Slattery, C. G. Wtndsor / Neutron diffraction study of stainless steel
176
A few examples of unidentifi~ peaks remained unaccounted for (table 2 and fig. 7). The most prominent examples seen in almost all specimens including the as-welded ones have d spacings around 5.80, 2.95 and 2.34 A. They may be the result of some slag inclusion since the d spacings do not correspond with any of the known phases given in table 5, the MC, carbide structure sometimes found in heat-treated stainless steel, nor to any of the higher-order reflections from the austenite peaks.
4. Discussion Type 316 weld metal possesses an inherently unstable microstructure due primarily to the presence of the metastable S-ferrite phase. This phase transforms fairly rapidly at the elevated temperatures encountered either by stress relieving the weld or by the thermal ageing taking place during in-service usage. It is important to understand the phase changes which take place and their subsequent implications for the mechanical behaviour of the weld. Essentially, the S-ferrite transforms to produce M,,C, carbide and intermetallic phases, primarily sigma and chi, together with the formation of additional austenite. The rate of transformation is relatively higher in the S-ferrite than in the austenite matrix since the S-ferrite is enriched locally in chromium and molybdenum such that the composition of the S-ferrite is close to the ranges over which sigma and chi exist as intermetallic phases. Consequently, on subsequent heat-treatment, the body-centred cubic ferrite can transform crystallographically with relative ease to the body-centred tetragonal sigma and body-centred cubic chi phases. The diffusion rate for intermetaIlic phase formation is higher in ferrite than in austenite and the
Table 8 A comparison transformation
of neutron diffraction and microscopical products after heat-treatments Initial ferrite content, as welded
Material
Low High Low High
ferrite ferrite ferrite ferrite
a C denotes
M,,C,
[22f on samples
Heat-treatment
BY mircroscopy
Neutron diffraction
(%) 6 18 6 18
(%) 6 16 6 16
carbide;
results
Cr/Mo segregation locally to the ferrite will reduce the mean path length over which diffusion must occur. In addition, the formation of M& carbide will occur even more readily since it involves rapid interstitial diffusion of carbon to the areas of local segregation of chromium. The appearance of any particular phase on thermal treatment depends on the specific temperature/time characteristics applied in the treatment and the corresponding behaviour patterns are normally described schematically by the appropriate time-temperatue-precipitation (mP) diagrams based on C-curve type kinetics induced by nucleation and diffusion-type reactions, [31]. This means that certain temperature ranges tend to favour particular phases, for example, at lower temperatures, M,,C, carbide is the preferred precipitate whilst at higher temperatures it can be sigma or chi intermetallics. Particular sequences of transformation will vary throughout the temperature ranges and the information obtained from a study of such tr~sformations will by its very nature be an average behaviour of the weld microstructure as a result of thermal treatment. In the present study, neutron diffraction techniques have been used to follow the phase transformations which occur in bulk samples of type 316 steel weldments. A similar study on the same welds using the more conventional techniques of optical and electron microscopy coupled with energy-dispersive analysis has shown very good agreement between the two methods and the two sets of results are compared in table 8. The presence of a high initial ferrite level in the weldment has a significant influence on the subsequent precipitation sequences during thermal treatment. Thus the 16% ferrite sample promotes the formation of the thermodynamically stable 1321 sigma phase after both the 85O*C and 62S’C treatments. Sigma-phase nucleation is re-
X the chi phase;
850°C/6 h 850°C/6 h 625”C/1050 h 625”C,‘lOOO h o the sigma phase;
taken
from
the same
Transformation
weldments:
products
nature
a
BY microscopy
Neutron diffraction
c + X + o(low)
c(1.5%)+s(l%) o(O.9%)+ S(l%) C( Is%)+ 6(4%) o(O.2%)+ S(4%)
0+-X c 4 6(4%) 0+X+6(6%)
and 6 the untransformed
ferrite.
of the
G. F. Slattety,
C. G. Wit&or
/ Neutron diffraction study of stainless steel
ported to be relatively easy at the higher temperature and its formation may occur without the prior appearance of M,,C, carbide [31]. The presence of sigma is usually considered to be undesirable since it may lead to reductions in long term creep ductility although its presence in this instance is below the 1% volume fraction level and may not be enough to affect overall properties. In contrast, the 6% ferrite sample produced M,,C, carbide as the main transformation product after both heat treatments and to the same volume fraction (1.5%) in both cases. This agrees well with the microscopic findings, particularly at 625°C and also with previous studies where carbide forms readily around this temperature region [10,33], by rapid diffusion of carbon to areas of local segregation of chromium and molydenum. M,& forms at the b-ferrite/austenite interface which represents a high angle boundary at which nucleation is favoured [31]. Even after ageing for 3000 h at 600°C, M,,C, carbide precipitates were still predominant and there was no evidence of sigma formation. By analogy with the precipitation sequence found in wrought 316 steel [15], the formation of M,,C, at 850°C is likely to be followed in the weld metal by the nucleation of sigma phase at the new S/y boundaries and growth of sigma should then continue with time. In this respect, the microscopic results were starting to show a small amount of sigma phase in the microstructure, table 8. A significant difference between the neutron diffraction and microscopical results is the absence of cm-phase in the neutron beam study, table 8. The appearance of chi-phase in the microscopical studies is particularly marked in the 850°C treatment. Lai and Haigh [lo] reported that chi-phase is the predominant precipitate in specimens heat treated at 800°C and is associated with the suppression of carbide. Chi forms by continuous transformation from the S/v interfaces and its formation is assisted by the prior presence of M,,C, carbide at these interfaces in the original “as-welded” structure [ll]. During the transformation of &ferrite, the product formed, particularly with regard to intermetallics is very dependent on variations in the composition of the weld metal together with its associated chromium equivalent [5]. The differences with respect to Chi arising from the two methods of examination may also reflect differences between bulk (neutron beam) and surface (SEM) examination effects in the weld specimen. During the transformation of S-ferrite, the amount of transformation products formed and hence the amount of ferrite remaining untransformed, depends very much on the precise composition of the initial weld metal [S].
111
The transformation of S-ferrite is reported to be rapid at temperatures above 750°C [33] although the present neutron study shows that about 1% ferrite is still retained untransformed after the 850°C/6 h treatment. Conversely, at lower temperatures, the transformation is reported to take place more slowly and this is reflected in the presence of 4% ferrite in the weld sample at the 625”C/lOOO h treatment. The agreement in the amount of retained ferrite between the microscopic and neutron beam methods is particularly good for this lower temperature treatment, table 8. The presence of b-ferrite remaining in the weld is of interest since S-ferrite is known to have a deleterious effect on the creep rupture strength of type 316 steel. During creep, microcracks can develop readily at the S-ferrite/austenite interphase boundaries [34]. The microcracks can then link into macrocracks provided there is sufficient S-ferrite present in the microstructure [35] but as in the case of the sigma phase, the presence of S-ferrite at the 1% (850°C treatment) and the 4% (625’C treatment) levels is probably insufficient to affect adversely the overall creep properties of the weldment. The associated scanning electron microscopical investigation [22] of these heat-treated weld specimens had shown that the chemical composition of the initial ferrite in the as-welded sample lay in the region of the ternary phase diagram for this alloy system which corresponded to the duplex ferrite and sigma-phase field. This was invoked to explain the readiness of the ferrite to transform to sigma phase on ageing since the ferrite already possessed the localised composition to transform easily to sigma, provided the crystallographic change from body-centred cubic to body-centred tetragonal could be accommodated. However, after ageing, the untransformed ferrite was found to be depleted in chromium and molybdenum owing to the formation of carbide and intermetallic phases. The resulting composition of this retained ferrite was found to have moved closer to that of austenite on the ternary diagram such that further transformation to sigma appeared unlikely or at least more difficult. This could explain the presence of a small but persistent amount of residual ferrite found after heat-treatment in the present neutron beam study.
5. Conclusions (a) Neutron diffraction techniques have been employed successfully to characterise the microstructures of Type 316 weld metal both as-welded and after heattreatment designed to simulate stress-relieving and inservice usage. The results compare well with previous
178
G.F. Slattety,
C.G. Windsor / Neutron diffraction study of stainless steel
findings utilising the more conventional techniques of optical and electron microscopy. Owing to their high penetration in steels, neutron beams have the great advantage, over microscopical techniques, of providing data characteristic of the bulk material. (b) Weld metal of two distinct ferrite levels were chosen for study. The “high’‘-ferrite weld metal of composition 20Cr 1ONi 3Mo contained 16% ferrite aswelded which on heat-treatment at 850°C for 6 h transformed to sigma phase of 0.9% volume fraction leaving 1% of ferrite untransformed. After 625’C for 1000 h the sigma phase formed was reduced in amount to 0.2% with 4% of untransformed ferrite. (c) The “low-ferrite weld metal of composition 18Cr 9Ni 2Mo contained 6% ferrite as-welded, which after 85O’C for 6 h transformed to Ma& carbide with a 1.5% volume fraction and 1% retained ferrite. After 625°C for 1000 h, the carbide produced also had a 1.5% volume fraction but with 4% of retained ferrite. (d) Neutron diffraction can provide important crystallographic data on the phases present in bulk weld specimens. For example, changes in lattice parameter of both the austenite matrix and the ferrite present in the weld can be followed after heat-treatment. Thus the austenite lattice was found to contract on both the S50°C/6 h and the 625’C/lOOO h treatments. The ferrite lattice contracted appreciably more than that of the austenite, especially for the 625”C/lOOO h treatment.
References
[l] F.C. Hull, Welding J. Res. Supplement 46 (1967) 3993. [2] J.C. Borland and R.N. Younger, Brit. Welding J. 7 (1960) 22. [3] Y. Arata, F. Matsuda and S. Katayama, Trans. Japanese Welding Res. Institute 5 (1976) 35. [4] T. Boniszewski, Metals and Materials (December 1978/January 1979) 41. (51 C.A.P. Horton, P. Marshall and R.G. Thomas, CERL Note No RD/L/2045 N81: Also Int. Conf. on Mech. Behaviour and Nuclear Applications of Stainless Steel at Elevated Temps., Varese, Italy (May 1981). [6] G.J. Curtis and N. Ibrahim, Metal Sci. 15 (1981) 566. [7] B.L. Baikie, A.R. Wagg, M.J. Whittle and D. Yapp, J. Brit, Nucl. En. Sot. 15 (1976) 257.
[8] S.R. Keown 386.
and R.G. Thomas,
Metal Science,
15 (1981)
[9] F.R. Beckitt, J. Iron and Steel Inst. 207 (1969) 632. [lo] J.K. Lai and J.R. Haigh, Welding J. Res. Supplement 58 (1979) 15. [ll] G.F. Slattery and P. G’Riordan, Metallography 13 (1980) 59. [12] H. Astrom, B. Loberg, B. Bergtsson and K.E. Easterling. Metal Sci. 10 (1976) 225. [13] C.E. Lyman, P.E. Manning, D.J. Duquette and E. Hall, Scanning El. Micr. 1 (1978) 213. [14] C.E. Lyman, Welding J. Res. Supplement (July 1979) 1895. [ 151 B. Weiss and R. Stickler, Trans. Met. Sot. AIME. 3 (1972) 851. [ 161 J.E. Spruiell, J.A. Scott, C.S. Airy and R.L. Hardin. Trans. Met. Sot. AIME. 4 (1973) 1533. [17] W.E. White and I. LeMay, Metallography 5 (1972) 333. [18] J.K.L. Lai and M. Meshkat, Metal Sci. (September 1978) 415. [19] J.K.L. Lai, D.J. Chaste11 and P.E.J. Flewitt. Mat. Sci. and Eng. 49 (1981) 19. [20] L.P. Stoter, J. Mat. Sci. 16 (1981) 1039. [21] G.F. Slattery, P. O’Riordan, M.E. Lambert and S.M. Green, J. Microscopy 122 Pt. 2 (1981) 109. [22] G.F. Slattery, S.R. Keown and M.E. Lambert UKAEA Report ND-R-829(R) 1982; also to be published in Metals Technology Journal. [23] C. Windsor and C. Wright, New Scientist, Vol. 88, No. 1231 (Dee) 714. 1241 S.A. David, Welding J. Res. Supplement, April (1981) 63s. [25] N. Suutala, T. Takalo and T. Moisio, Trans. Met. Sot. AIME 10A (1979) 512. [26] A.H. Baston and D.H.C. Harris, eds., AERE Rep. No. 9278 (1978). [27] M. Sakata, R. Sinclair and C. Windsor, unpublished work ( 1980). [28] G.E. Bacon, Neutron ford, 1975) p. 96.
Diffraction
(Clarendon
Press, Ox-
[29] M.J. Cooper, K.D. Rouse and R. Sayers, unpublished work 1977. [30] M.J. Cooper and K.D. Rouse, J. Appl. Cryst. 14 (1981) 72. [31] R.G. Thomas and S.R. Keown, CEGB Report RD/M/R312. [32] C.J. Bechtoldt and A.C. Vacher, J. Res. Nat. Bur. Stand. 58 (1957) 7. [33] R.G. Thomas and D. Yapp, CEGB Report R/M/R273. [34] I. LeMay, B.J. Bassett and W.E. White, Paper Cl78 in: Creep and Fatigue in Elevated Temp. Applications (Inst. Mech. Engrs., London, 1974). [35] J.R. Haigh, Welding J. Res. supplement, 56(5) (1977) 149s.
Journal of Nuclear Materials 118 (1983) 165-178 North-Holland Publishing Company
THE APPLICATION OF NEUTRON DIFFRACTION TO A STUDY OF PHASES IN TYPE 316 STAINLESS STEEL WELD METALS G.F.
SLATT’ERY
UKAEA,
Risley Nuclear Power Development
Laboratories,
Risiey, Warrington
WA3 6AT, UK
and C.G.
WINDSOR
UKAEA, Materials Physics Division, Atomic Energy Research Establishment, Received
21 April
1983; accepted
Hanvell, Oxon, OX1 I ORA, UK
6 June 1983
Neutron diffraction techniques have been utilised to study the phases in type 316 austenitic stainless steel weld metal, both in the as-welded condition and after stress-relieving and ageing heat-treatments. The amounts of the principal crystallographic phases present in bulk specimens have been measured. Two compositions of weld metal were selected to provide a “low” (6%) and “high” (16%) initial ferrite level and the subsequent volume fractions of transformation products were measured after heat-treatment. Some retained ferrite was observed in all the heat-treated specimens, ranging from 4% for specimens of both initial ferrite levels treated at 625’C for 1000 h, to around 1X for the specimens treated at 85O’C for 6 h. The high initial ferrite specimen produced 0.9% of sigma phase after the 850°C treatment and 0.2% sigma after the 625°C treatment. The low initial ferrite specimen produced 1.5% M,,C, carbide after both heat-treatments. The results compare well with previous findings on similar samples of weld metal using optical and electron microscopy.
1. Introduction Welds in type 316 austenitic stainless steel are of considerable interest because of their usage in both conventional and nuclear power stations. As a result, a high standard of structural integrity is expected from the weldment. The composition of the weld is formulated so as to provide between about three and eight per cent of delta-ferrite in the microstructure to avoid hot tearing problems during weld solidification [l-4]. The amount of S-ferrite present is essentially dependent on the chemical composition of the weld deposit and various combinations of chromium, nickel and molybdenum are used in the weld metal to provide the optimum amount of S-ferrite. Typical compositions of the weld metal used in service are 19Cr12Ni3Mo and the leaner 17Cr8Ni2Mo [5]. In practice, the weld metal is deposited as a succession of overlaid weld beads. The macrostructure of such multipass weld layers in influenced by the absence of any phase transformation in the weld on cooling, with the result that the structure consists of long, columnar
0022-3 115/83/0000-0000/$03.00
0 1983 North-Holland
grains which nucleate from favourably oriented grains in the heat-affected zone and grow epitaxially from bead to bead in the weld [5]. The columnar grains grow along lines of heat dissipation such that a strong elastically anisotropic character is developed in the weld metal. This feature of marked anisotropy renders the ultrasonic inspection of such welds difficult owing to directional variability of the velocity and attenuation of the elastic waves [6,7]. Each grain contains a network of S-ferrite distributed throughout the austenite matrix such that the fine-scale structures consists of a series of low-angle boundaries or dislocation networks linking the S-ferrite into a sub-grain structure with a diameter of less than 10 pm [8]. Large components of weld metal need to be stressrelieved for service usage to reduce or even eliminate the residual stresses introduced during the welding cycle. During such heat-treatment, the S-ferrite phase in the weld microstructure transforms to produce M&, carbide and intermetallic phases such as chi and sigma [8-111. Since the S-ferrite is enriched in chromium and molybdenum due to microsegregational effects [ 12-141
166
G.F. Slattery C.G. Windsor / Neutron drffaction study of starnless sreel
then these transformations occur much more rapidly in the ferrite phase or at the ferrite-austenite interface than in the austenite matrix. Consequently, these reactions take place more readily in a duplex weld structure than in the equivalent wholly austenitic parent steel. Since the production of these phases on heat-treatment affects both the chemical and mechanical properties of the weld metal, it is important to understand how these properties may be modified by changes in composition and microstructure of the weld metal. Thus changes in composition of the weld metal can affect not only the amount of S-ferrite present in the original weld but also the amount of intermetallic phases formed as a result of the transformation of the S-ferrite during subsequent heat-treatment [5]. A study was initiated to investigate the influence of different levels of a-ferrite on the phase transformations produced as a result of both a stress relieving treatment at 85O’C and also an ageing treatment at 625°C to simulate an in-service usage. Since previous work has been carried out on the structure of type 316 steel both in the wrought and welded conditions using primarily electron-microscopical techniques [ 12-201, a detailed study was first made on these stress-relieved and aged welds using techniques of microanalysis with the electron microscope together with selective etching methods and optical microscopy [21]. This work is being reported separately [22]. The second stage of the work reported here, utilises the alternative technique of neutron-beam diffraction which is a complementary technique to more conventional X-ray diffraction. The study assesses the feasibility of using neutron diffraction on bulk samples of significant thickness to provide the bulk volume fractions of the phases in the welds in order to correlate the results with the previous electron/optical microscopic studies on the identical samples of weld material. Neutron diffraction has the great advantage that the beam is able to penetrate metals since, having no electric charge or electric field, it interacts only weakly with matter [23]. Neutron beams can therefore be used to study bulk effects for example in steels, the penetration depth is approximately 8 mm. Also the surface condition of the sample is unimportant so that no special preparation techniques are needed. In contrast, electron beams interact much more easily and so penetrate essentially only comparatively small distances. In the present study, neutron diffraction was used (1) to measure the volume fraction of S-ferrite originally present in the weld samples together with the amount of untransformed S-ferrite after heat treatment, (2) to obtain basic crystallographic information on the face-centred cubic y-austenite and body-centred cubic S-ferrite phases and
(3) to identify and quantify the amount of the phases produced as a result of the transformation of d-ferrite during heat treatment.
2. Experimental procedure 2.1. Materials and heat-treatment Two commercially available type 316 welding electrodes to B.S. 2926: 1970 were chosen to give approximately 5 and 15% of b-ferrite in the weld metal. Their compositions and actual ferrite levels are given in table 1 which shows that the welds had nominal compositions lKr9Ni2Mo and 20Cr lONi3Mo. The increased levels of the ferrite-forming chromium and molybdenum in the second composition resulted in the higher level of d-ferrite content. The weld deposits were laid down by manual metallic arc-welding to give weld pads 150 x 25mm in size typical welding conditions being: current, 125 A; voltage, 20 V; electrode size, 4 mm diameter; heat input, 1 kJ; with maximum interpass temperature, 150°C. Specimens were examined in the neutron beam diffractometer in the as-welded condition and also after the following heat treatments, namely (1) stress-relieved at 850°C for 6 h and (2) aged at 625°C for 1000 h. Microscopical examination showed that as-welded the “low” ferrite steel (6%) possesses a discontinuous network of ferrite with a vermicular type morphology [24,25] as the minor phase distributed within an austenite matrix. In the “high” ferrite weld (16%), the a-ferrite formed a more continuous network characteristic of the morphology designated as “lacy” in type [24]. Also present in the general microstructures were small spheroidal slag particles approximately 0.5 to 3 pm in
Table 1 Composition (wtW) and ferrite content of welds Element
“High” ferrite weld metal
“Low” ferrite weld metal
Carbon Manganese Silicon Chromium Nickel Molybdenum Phosphorus Sulphur Ferrite volume fraction (%)
0.047 0.96 0.41 20.5 10.1 3.13 0.02 0.01
0.067 2.18 0.35 18.0 8.9 1.85 0.02 0.01
15.6
5.9
167
G.F. Slattery, C.G. Windsor / Neutron diffraction study of stainless steel
Fig. 1. Scanning electron micrograph of a heat treated d-ferrite network of the duplex weld metal.
weld sample showing
precipitation
of sigma and chi intermetallic COUNTER
phases in the
SHIELO I
size. Fig. 1 is a scanning electron microscope image of a heat-treated weld sample showing precipitation of sigma and chi intermetallic phases in the S-ferrite network of the duplex weld metal. 2.2. Neutron
MONOCHROMATOR
diffraction
The experimental technique is very similar to X-ray diffraction with the substitution of the nuclear reactor for the X-ray source. Neutrons of a particular wavelength X are selected from the Maxwellian distribution of neutron wavelengths present in the core by Bragg reflection from a single crystal. The beam direction is defined by means of a slit collimator, and the beam size adjusted so that it matches the sample size. The scattered neutrons are detected by one or more gas counters. The scattered neutron intensity is measured as a function of
-_____ Fig. 2. The layout reactor, Harwell.
~of the CURRAN
diffractometer
at DID0
BEAM
STOP
v
G.F. Slatrery, C.G. Windsor / Neutron diffraction study of stainless steel
168
28 = 90”. The instrumental peak width was fitted to the formula W= A sin 8/AB with A = 806”/nm and 3 =
the scattering angle “A” to plot out the diffraction pattern. Fig. 2. shows the layout of the CURRAN diffractometer in the DID0 reactor at Harwell used in these measurements 1261.The monochromator used was hot pressed “squashed” germanium at a take-off angle 26, = 47.5’. Using the (331) germanium reflection gave a wavelength of (1.369 f 0.002) A as determined from a nickel powder calibration scan. This wavelength is convenient in that the second-order scattering (h/2) from the mon~hromator is negligible, while the range of d spacings down to about 0.1 nm covered by the scan covers several austenitic and ferritic peaks. The counter bank contains 5 distinct counters with approximately 10” separation. The exact separation and the nominal zero of the scattering angle is found from the nickel powder scan. The spectrometer was used with a collimation of 0.5” before the sample and about lo geometric collimation after the sample. This gave peak widths of order 0.4’ giving a resoiution of Ad/d = 0.008 near 28 = 4S”, rising to a width of 1.5Oor Ad/d = 0.014 near
RIS7
2.41. Relative intensity between the counters was calibrated by means of a previous calibration run using a vanadium incoherent scattering sample. 2.3. Sample preparation The samples for diffraction examination were sawn from the weld pad. They were approximately 27 mm high, 15 mm wide and 2 mm thick. They were mounted on the spectrometer with the base of the sample pressed into a piece of plasticene, covered in cadmium foil to act as a neutron absorber. This was in turn mounted on a base which rotated at about 4 revolutions per minute. This rotation was made in order to reduce preferred orientation effects in the sample. Since these were not eliminated by rotation in only one axis, the results will still contain an inaccuracy from residual preferred orientation effects. The sample was centred in the neutron beam by taking photographs using a scintillator plate holder.
HIGH
6H850
6O.OC
0.00
5.00
1
I
15.00
25.00
35.00
45.00 2-THETA
Fig.
3.
I
,
,
55.00
65.00
75.00
fOEGREEt
An example of the raw diffraction pattern on a scale to include the matrix peaks.
L
85.00
1
95.00
G. F. Slattety,
C.G. Windsor / Neutron diffraction study of stainless steel
r: c I
t
RIS
11
I
i
IGH
6HE.O
fig. 4. The raw neutron diffraction patterns on an expanded scale. The zeros of the graph are suppressed and no normalisation has been performed. Most of the runs lasted about 20 h and corresponded to a scan of from 7” to 102” in 0.2’ steps. The counting time at each point was varied to correspond to fixed
counts on a fission chamber monitor counter placed in the incident beam. By this means, variations in the incident flux are cancelled out. A separate run was
G.F. Slattery, C.G. Windsor / Neutron diffvuction study
170
made without a sample in order to determine the instrumental background. This turned out to be small compared to the disorder scattering from the several components of these steels together with incoherent scattering from the sample at ali angles. A specimen of nickel powder was mounted and scanned in the same way as
Table 2 Interplanar Probable assignment
d spacings
of observed
reflections
the other specimens in order to determine zero angle and the wavelength. 2.4. Data processing
Data from each run were punched on paper tape which was analysed on a DEC PDP 1 l/40 minicom-
1000 h at 625°C
6 h at 850°C
High
Low
High
Low
1.0183(g)
1.0844(5)
1.9394(5) 1.0857(2)
1.O 124(30) 1.0379(5) 1.0841(2)
1.0379(5) 1.0851(2)
I .0380(4) 1.0841(27)
a 211 (1, 720 y 220
1.1742(6)
l.i780(3)
1.1712(7)
1.1728(3)
I. 1728(7)
1.2770(2)
1.273415)
1.2717(4)
1.2727(2)
1.2717(2)
cr c
200 640
1.4380(4)
1.4380(2)
1.4348(28) 1.4812(42)
1.4364(7)
1.4367(12) 1.4842( 17)
c d (I c7
622 501 322 312
Y y 0
200 222 331
cy 220 y 222 y 311
1.03&l(8)
1.5627( 17) 1.6111(30)
0 a y G c d
1.4927(U) 1.5505(6)
I .6066(45) 1.6388
1.6485(6) 1.6742(3)
I .7980(2)
I .8007(2)
1.7995(2)
320
rJ
310
c 0 c
220 101 200
1.7980(2)
1.7992(2) 1.8455(16) I .8931(4)
1.8893(17)
2.0340(5) 2.0766(3) 2.1267( 180)
2.0364(5) 2.0787(3)
2.0300( 10) 2.0761(3)
I .88 I7(59) 1.9373(4) 1.9827(5) 2.0315(5) 2.0776(3) 2.1327(27)
1.9377(4) 1.9841(5) 2.0330( 10) 2.0756(3)
2.0364(5) 2.0771(3) 2.1404(11)
2.1753(29) 2.3423(67)
0
1.0094(2) 1.0388( 1) I .0848(2) 1.1199(5) 1.1481(8) 1.1764(6) 1.2392(33) I .2724(2) 1.3315(6) 1.3837(15) 1.4360(9)
1.6060(60)
1.8813(13) 411 212 110 111 410 422 002
High
1.7659(22) 1.7988(4)
C440 0
the counter
(A) a
As welded Low
of starnless steel
2.3423(67)
2.3464(134) 2.6806(35)
2.9528(2 16)
2.9474(32)
2.9495(86) 3.7165(87) 5.28
1l(600)
5.7794(127)
2.2902( 19) 2.3551(20) 2.4447(7) 2.6648( 123)
2.3332(66) 2.3330(66)
2.2966( 13) 2.3558(20)
2.9593(43)
2.9420(2 15) 3.7287( 192) 4.1396(43) 5.3059(36) 5.8006(980)
2.6771(9) 2.7701(29) 2.9528(43) 3.6771(51) 4.0635(21) 5.2357( 173) 5.8609(44)
a The numbers in brackets refer to the statistical variance of the least squares fit to the d spacing. Absolute of 0.02% because
of systematic
errors
in the wavelength
calibration.
errors are larger by order
171
G.F. Slattety, C.G. Windsor / Neutron diffraction study of stainless steel
From the fitted peak scattering angles, the d spacing of each peak was calculated from the Bragg equation. The d spacings of all the located peaks are given in table 2 together with an estimate of the statistical error. In nearly all cases, the width of the reflection was slightly greater than the instrumental width given by the equation
puter. The first program PATN4 [27] takes the data from all the six counters and combines them together into a single file corrected for instrumental background zero angle shifts, counter efficiencies, and counter separations. Fig. 3 shows the peaks from the austenitic matrix on a scale to fit the graph. Fig. 4 shows the results on a greatly expanded scale to emphasise the minority phase peaks. A second program ROBOT locates the peaks automatically and finds their appropriate position and intensity. The third program CRAWL takes this file and refines the peak intensities, positions and widths over and above the resolution. The program also gives an estimate of the statistical errors in these variables. FERRITE 200
T
5 the
the
extra widths are shown above widths and expressed in the form
cam
8628
2~
2dtane
I
OD
//
628
(3)
.
A slope on the figure corresponds to strain broadening. It is such that both the austenite and ferrite peaks have a crystallite size of around 60 nm but that the ferrite peaks appear to have some additional strain broadening.
1
AS WELDED LOW FERRITE
(2)
’
2x
where A 28 is the extra width of the angular scan. In this figure, a flat line corresponds to a crystallite size.
500
I
fig.
L=Q=
I
0
(1)
AQ -= 4s
I 211
__+__'---I--
t
8,,wn/2.
In
_---
0 006
COS
instrumental
PEAKS
I
i 110
D= d
4
3. Results 3.1. Analysis
--
0
f
1000 HOURS AT 625'C LOW FERRITE I
ce ,
/
.* -
of the austenite and ferrite lattice constants
Austenite has a face-centred cubic structure (fee). For cubic structures, the lattice constant can be derived for any peak of known Miller indices (hkl) from the equation a, = d( h2 + k2 + 12)“2.
0
0
I I
I
;K$RS FERRITE
AT .350°C 111 200 AUSTENITE 1
1
'I
f
6 HOURS AT 850°C HIGH FERRITE
,PEAKS 0.1 1 Sin
Table 3 Lattice constants and volume phases in weld metal
B/h
f 220 1
1 311 222 I 0.2
IOD
,2r=
Fig. 5. The line widths of the austenite and ferrite reflections, above the resolution value, expressed as a broadening factor Q (left hand scale) or as a crystallite size (right hand scale). The finite slope of the ferrite peak widths corresponds to some strain broadening.
Austenite “Low” ferrite “High” ferrite Ferrite “ Low” ferrite “High” ferrite Ferrite volume “ Low” ferrite “ High” ferrite
(4)
fraction
of austenite
and ferrite
As welded
1OOOh at 625°C
6h at 850°C
3.5972(5) 3.601 l(6)
3.5959(6) 3.5991(5)
3.5959(7) 3.5982(4)
2.8763(2) 2.8792(2) fraction (%) 5.9 15.6
2.8682(5) 2.8729(S)
2.8737( 12) 2.8778(51)
4.0 3.5
1.6 1.0
172
G.F. Slattery,C.G. Windsor / Neutron diffracrion study of stainless steel 3.2. Volume fractions of the ferrite phase Neutron scattering intensities depend on many factors but all of these can be readily evaluated so that absolute intensity measurements are straightforward to an accuracy of a few per cent (to achieve the highest accuracy careful corrections must be made for factors such as resolution, multiple scattering and extinction). In the present context, it is sufficient to obtain intensities relative to the austenitic matrix so that many factors cancel. The total intensity of a powder peak integrated over the angular scan is given by Bacon as follows [28]
AUSTENITE
1 NZ,,,[F,,,12X3L,exp(-2W)
(I=-
""7t
, LOW AS WELDED
V sin Bra sin 20
8772
HIGH FERRITE-t
FERRITE-! lOOOh 625%
FERRITE 6h 850°C
Fig. 6. The lattice constants of the austenite and ferrite phases in “high” and “low” ferrite specimens as welded and after heat treatments.
’
assuming N unit cells each of volume V and where Z,,, is the multiplicity of reflection h/cl. The Lorentz factor, LJ2rsin28, represents the fraction of the whole Debye-Scherrer cone intercepted by the counter. The last factor is the Debye-Waller factor representing thermal motions. Fhk, is the structure factor depending on the mean neutron scattering length b of atoms in positions R in the unit cell such that Fhk,=bexp(-iQR).
(6)
For a primitive structure cell, such as the austenite The five reflections 111, 200, 220, 311 and 222 corresponding to the austenite pattern were always clearly visible, and all gave consistent values of the lattice constant. The values given in table 3 and plotted in fig. 6 show the mean lattice constant, and its variance determined from these five reflections. In the case of the body-centred cubic (bee) ferrite peaks, the accuracy is somewhat less. Only three reflections 110,200, and 211 were covered by the range of the experiment. Of these the alpha (110) peak was only seen as a “wing” to the gamma (111) peak, and was also unresolved from the sigma (202) reflection. Also the alpha (220) was very close to the gamma (222) reflection. Thus in practice the alpha (110) peak could not be used in sigma-containing steels and only two or three alpha reflections were used in determining the ferrite lattice constant. The results in fig. 6 show a clear and significant change in lattice parameter with heat-treatment. The austenite lattice contracts on both the 850°C/6 h and 625O/lOOO h heat-treatments. The ferrite lattice shrinks appreciably, more especially for the 625°C/1000 h treatment. However the “high” ferrite (16%) lattice constants are consistently more expanded than the “low” ferrite (6%) lattice constants.
4~
with only one atom per unit and ferrite phases,
= b
and thus (7 sin 0 sin 28 Z,,,
NA3b2
exp( - 2W) = F
(7)
= ”
that is, a constant for any given set of reflections for a particular phase. The relative number of unit cells for two phases can therefore be found from the ratio of the constants C. For example, for the ratio of body-centred cubic to face-centred cubic phases, we have
N
Vbccb:,, Cbcc
N FCC
vfcc
bee -=---
b;_
Table 4 Neutron scattering
Cfcc .
cross-sections
Element
At. wt
b (IO-”
Fe Cr Ni MO Si C
55.8 52.0 58.7 95.9 28.06 12
0.95 0.352 I .03 0.69 0.42 0.665
cm)
173
G. F. Slattery, C.G. Windsor / Neutron diffrraction study of stainless steel Table 5 The matrix and secondary phases in type 3 16 steel weld metal: approximate compositions
Phase
Crystal structure
Austenite Delta-ferrite Ma sCg carbide Sigma Chi
face centred cubic body centred cubic cubic tetragonal cubic
AI A2 A2 D8b A12
Fe (wt%)
Cr
Ni
Ma
68 66 2s 58 57
18 23 65 28 23
9 5
2 4 5 8 13
The neutron scaitering iengths for the nuclei present in type 316 stainless steel are given in tabfe 4. The phases identified in this work are listed in table 5 and using the atomic concentrations measured for the austenite and ferrite phases given in this table, the mean scattering lengths can be computed as b kc = 7.81 X b fee =
IO-“mm
8.47 X lo- I2 mm
The Debye-Waher factors for pure iron and for type 316 stainless steel have been determined by previous diffraction studies to high values of sin e/X on the
LINAC pulsed neutron fusing the conventional Wailer factors) are B roc=0.47A2;
source. Appropriate values symbohsm for the Debye-
B,,,=0.29$.
The constant C was computed for each reflection and a mean taken. Table 6 shows experimental and calcuhrted line intensities for the austenite and ferrite reflections. It is seen from table 6 and from fig. 7 where the experimental and calculated line intensities are plotted on a logarithmic scale against angle, that the intensities are genera& in agreement but that the calculated and observed intensities are occasionally in
Table 6 Caktdated and observed line intensities for the austenite and ferrite f& refieetions Reflections (W
Materiai As welded
Austenite 111 200 220 311 222 Ferrite
110 200 211
ObLWNd
calculated observed caktiated observed calculated observed calculated observed calculated observed calculated observed calculated observed calculated
la00 h at 625OC
6 h at 85O’C
Low %
High 6
Low 6
High 6
Low 6
High s
53 46 48 26 20 26 24 39 4.5 12.2 43 32 9.8 8.4 II 23
48 51 71 29 20 29 25 44 3.6 1.4 122 96 35 25 33 69
49 34 25 19 14 19 18 29 4.5 8.9 22 16 3.5 4.1 6.1 11.5
29
84 76 86 43 30 44 42 65 6.2 20.2 21 15 3.8 3.9 5.0 IO.8
68 73 101 41 20 42 36 63 5.4 19.4 12.7 8.6 0.62 2.22 7.8 6.2
29 36 16 13 17 15 25 2.5 7.7 28 I2 3.5 3.1 8.3 8.7
G. F. Slattery, C.G. Windsor / Neutron diffraction study of stainless steel
174
cause of the systematic error resulting from preferred orientation. The results in table 3 show that the as-welded specimens have about 6 and 16% ferrite as the duplex constituent with austenite. With either original ferrite percentage, around 4% ferrite remains after the ageing treatment at 625°C for 1000 h and 1% remains after the stress-relieving treatment at 850°C for 6 h. 3.3. The identification of precipitate phases
1000
c
LOW
The phases identified by electron microscopical techniques in type 3 16 stainless steel weld metal [ 10,221 and listed in table 5 are all of known structure [ 151 so that it was possible to predict the approximate d spacings and relative peak intensities expected from each structure. The d spacings were calculated with the KDRREF programme [29] and the line intensities by the TAILS programme [30]. Fig. 8 shows the calculated scattering angles and intensities on a logarithmic scale and to the same scale as the experimental results in fig 7. By laying one over the other, it is straightforward to identify those phases having several peaks in the observed spectrum.
HOURS 625Y FERRITE
3
L
HIGH
FERRITE
3
4 3 2 1 10
20
30
LO
50
60
2 80
Fig. 7. Experimental and calculated line intensities for the reflections plotted on a logarithmic scale versus angle. The located peaks in the six specimens are indicated by solid lines. The dashed line and side mark indicate calculated intensities.
error by a factor of two or three. This is probably the result of the appreciable preferred orientation which was present in the sample. The average value over a number of reflections which is used in determining the volume fraction should be appreciably more accurate. Table 3, lower section, shows the volume fraction of ferrite compared to austenite evaluated using eq. (8). The number is strictly a number fraction N,,,/N,,, but since the atomic volumes are so nearly equal, it is equal to the number fraction within experimental error. The statistical error on the volume fraction is negligible, but an overall error of perhaps 20% must be allowed be-
-PHASE 2I
1
3
Y
I
I
1
PHASE
2 II
100
I
200
I
300
I
1
400
500
60"
29s
Fig. 8. The calculated intensities and scattering angles of the reflections from phases known to be present in the steels.
175
G. F. Slattery, C. G. Windsor / Neutron diffraction study of stainless steel
By this means, the sigma phase could be clearly identified in the “high” ferrite specimens heat-treated at both 85O’C and 625°C. It was also clear that the M,,C, carbide structure was present in the “low” ferrite specimens heat-treated at both temperatures. However no other phases could be identified in the observed patterns and none of the minority precipitate phases such as carbide and intermetallics could be identified in the as-welded specimens. The sigma-phase lattice parameters were deduced from the 002 and 410 reflections and the carbide phase parameter from the 440 reflection as given in table 7. The line intensities calculated with the TAILS program were then matched with the experimental line intensities in the cases given in the table. The
total intensity summed over the lines given were then used to calculate the volume fraction of the phase concerned. Once again the systematic error caused by preferred orientation leads to an error of the order of 20% in these volume fractions. In all cases, the volume fraction has been evaluated by multiplying the relative number of unit cells obtained from the TAILS intensities by the number of atoms per unit cell for each phase. In the “high” ferrite sample after the 850°C heat-treatment, the sigma phase volume fraction was 0.9% but in the sample with the 625’C treatment it was only 0.2%. In the “low” ferrite samples, the volume fraction of carbide after both heat-treatments was about 1.5%.
Table 7 A comparison of the observed intensities (upper figures) with calculated carbide and sigma phases in the heat-treated weld specimens
intensities
(lower figures)
for different
reflections
from the
Sigma phase high-ferrite specimen
Carbide (M,,C,) phase low-ferrite specimen Treatment:
1000 h at 625OC
6 h at 850°C
Treatment:
1000 hat 625°C
6 h at 850°C
Lattice a, (A) parameters
10.64
10.64
a0 (A) CO
8.8 4.60
8.8 4.593
Int.
Int. 171 430 90 74 87 92
(W
Int.
Int.
(hW
200
(122) (20) (121) (34) 70 15 71 128 46 16 71 103 52 59
(151) (36) 126 62
101
222 400 440 533 622 711,640
Ill 310
135 231
320 002
113 186 77 106
410 212 411 331
85 13 126 122 195 186 163 118 259 392 186 182
222 312 322 501 Volume fractions
N( c)/N(
y)
0.015
0.015
NAN
0.0022
589 1217 1910 1850 1907 1176 3282 3893 2155 1818 935 347 604 477 224 324 142 389 0.0091
G.F. Slattery, C. G. Wtndsor / Neutron diffraction study of stainless steel
176
A few examples of unidentifi~ peaks remained unaccounted for (table 2 and fig. 7). The most prominent examples seen in almost all specimens including the as-welded ones have d spacings around 5.80, 2.95 and 2.34 A. They may be the result of some slag inclusion since the d spacings do not correspond with any of the known phases given in table 5, the MC, carbide structure sometimes found in heat-treated stainless steel, nor to any of the higher-order reflections from the austenite peaks.
4. Discussion Type 316 weld metal possesses an inherently unstable microstructure due primarily to the presence of the metastable S-ferrite phase. This phase transforms fairly rapidly at the elevated temperatures encountered either by stress relieving the weld or by the thermal ageing taking place during in-service usage. It is important to understand the phase changes which take place and their subsequent implications for the mechanical behaviour of the weld. Essentially, the S-ferrite transforms to produce M,,C, carbide and intermetallic phases, primarily sigma and chi, together with the formation of additional austenite. The rate of transformation is relatively higher in the S-ferrite than in the austenite matrix since the S-ferrite is enriched locally in chromium and molybdenum such that the composition of the S-ferrite is close to the ranges over which sigma and chi exist as intermetallic phases. Consequently, on subsequent heat-treatment, the body-centred cubic ferrite can transform crystallographically with relative ease to the body-centred tetragonal sigma and body-centred cubic chi phases. The diffusion rate for intermetaIlic phase formation is higher in ferrite than in austenite and the
Table 8 A comparison transformation
of neutron diffraction and microscopical products after heat-treatments Initial ferrite content, as welded
Material
Low High Low High
ferrite ferrite ferrite ferrite
a C denotes
M,,C,
[22f on samples
Heat-treatment
BY mircroscopy
Neutron diffraction
(%) 6 18 6 18
(%) 6 16 6 16
carbide;
results
Cr/Mo segregation locally to the ferrite will reduce the mean path length over which diffusion must occur. In addition, the formation of M& carbide will occur even more readily since it involves rapid interstitial diffusion of carbon to the areas of local segregation of chromium. The appearance of any particular phase on thermal treatment depends on the specific temperature/time characteristics applied in the treatment and the corresponding behaviour patterns are normally described schematically by the appropriate time-temperatue-precipitation (mP) diagrams based on C-curve type kinetics induced by nucleation and diffusion-type reactions, [31]. This means that certain temperature ranges tend to favour particular phases, for example, at lower temperatures, M,,C, carbide is the preferred precipitate whilst at higher temperatures it can be sigma or chi intermetallics. Particular sequences of transformation will vary throughout the temperature ranges and the information obtained from a study of such tr~sformations will by its very nature be an average behaviour of the weld microstructure as a result of thermal treatment. In the present study, neutron diffraction techniques have been used to follow the phase transformations which occur in bulk samples of type 316 steel weldments. A similar study on the same welds using the more conventional techniques of optical and electron microscopy coupled with energy-dispersive analysis has shown very good agreement between the two methods and the two sets of results are compared in table 8. The presence of a high initial ferrite level in the weldment has a significant influence on the subsequent precipitation sequences during thermal treatment. Thus the 16% ferrite sample promotes the formation of the thermodynamically stable 1321 sigma phase after both the 85O*C and 62S’C treatments. Sigma-phase nucleation is re-
X the chi phase;
850°C/6 h 850°C/6 h 625”C/1050 h 625”C,‘lOOO h o the sigma phase;
taken
from
the same
Transformation
weldments:
products
nature
a
BY microscopy
Neutron diffraction
c + X + o(low)
c(1.5%)+s(l%) o(O.9%)+ S(l%) C( Is%)+ 6(4%) o(O.2%)+ S(4%)
0+-X c 4 6(4%) 0+X+6(6%)
and 6 the untransformed
ferrite.
of the
G. F. Slattety,
C. G. Wit&or
/ Neutron diffraction study of stainless steel
ported to be relatively easy at the higher temperature and its formation may occur without the prior appearance of M,,C, carbide [31]. The presence of sigma is usually considered to be undesirable since it may lead to reductions in long term creep ductility although its presence in this instance is below the 1% volume fraction level and may not be enough to affect overall properties. In contrast, the 6% ferrite sample produced M,,C, carbide as the main transformation product after both heat treatments and to the same volume fraction (1.5%) in both cases. This agrees well with the microscopic findings, particularly at 625°C and also with previous studies where carbide forms readily around this temperature region [10,33], by rapid diffusion of carbon to areas of local segregation of chromium and molydenum. M,& forms at the b-ferrite/austenite interface which represents a high angle boundary at which nucleation is favoured [31]. Even after ageing for 3000 h at 600°C, M,,C, carbide precipitates were still predominant and there was no evidence of sigma formation. By analogy with the precipitation sequence found in wrought 316 steel [15], the formation of M,,C, at 850°C is likely to be followed in the weld metal by the nucleation of sigma phase at the new S/y boundaries and growth of sigma should then continue with time. In this respect, the microscopic results were starting to show a small amount of sigma phase in the microstructure, table 8. A significant difference between the neutron diffraction and microscopical results is the absence of cm-phase in the neutron beam study, table 8. The appearance of chi-phase in the microscopical studies is particularly marked in the 850°C treatment. Lai and Haigh [lo] reported that chi-phase is the predominant precipitate in specimens heat treated at 800°C and is associated with the suppression of carbide. Chi forms by continuous transformation from the S/v interfaces and its formation is assisted by the prior presence of M,,C, carbide at these interfaces in the original “as-welded” structure [ll]. During the transformation of &ferrite, the product formed, particularly with regard to intermetallics is very dependent on variations in the composition of the weld metal together with its associated chromium equivalent [5]. The differences with respect to Chi arising from the two methods of examination may also reflect differences between bulk (neutron beam) and surface (SEM) examination effects in the weld specimen. During the transformation of S-ferrite, the amount of transformation products formed and hence the amount of ferrite remaining untransformed, depends very much on the precise composition of the initial weld metal [S].
111
The transformation of S-ferrite is reported to be rapid at temperatures above 750°C [33] although the present neutron study shows that about 1% ferrite is still retained untransformed after the 850°C/6 h treatment. Conversely, at lower temperatures, the transformation is reported to take place more slowly and this is reflected in the presence of 4% ferrite in the weld sample at the 625”C/lOOO h treatment. The agreement in the amount of retained ferrite between the microscopic and neutron beam methods is particularly good for this lower temperature treatment, table 8. The presence of b-ferrite remaining in the weld is of interest since S-ferrite is known to have a deleterious effect on the creep rupture strength of type 316 steel. During creep, microcracks can develop readily at the S-ferrite/austenite interphase boundaries [34]. The microcracks can then link into macrocracks provided there is sufficient S-ferrite present in the microstructure [35] but as in the case of the sigma phase, the presence of S-ferrite at the 1% (850°C treatment) and the 4% (625’C treatment) levels is probably insufficient to affect adversely the overall creep properties of the weldment. The associated scanning electron microscopical investigation [22] of these heat-treated weld specimens had shown that the chemical composition of the initial ferrite in the as-welded sample lay in the region of the ternary phase diagram for this alloy system which corresponded to the duplex ferrite and sigma-phase field. This was invoked to explain the readiness of the ferrite to transform to sigma phase on ageing since the ferrite already possessed the localised composition to transform easily to sigma, provided the crystallographic change from body-centred cubic to body-centred tetragonal could be accommodated. However, after ageing, the untransformed ferrite was found to be depleted in chromium and molybdenum owing to the formation of carbide and intermetallic phases. The resulting composition of this retained ferrite was found to have moved closer to that of austenite on the ternary diagram such that further transformation to sigma appeared unlikely or at least more difficult. This could explain the presence of a small but persistent amount of residual ferrite found after heat-treatment in the present neutron beam study.
5. Conclusions (a) Neutron diffraction techniques have been employed successfully to characterise the microstructures of Type 316 weld metal both as-welded and after heattreatment designed to simulate stress-relieving and inservice usage. The results compare well with previous
178
G.F. Slattety,
C.G. Windsor / Neutron diffraction study of stainless steel
findings utilising the more conventional techniques of optical and electron microscopy. Owing to their high penetration in steels, neutron beams have the great advantage, over microscopical techniques, of providing data characteristic of the bulk material. (b) Weld metal of two distinct ferrite levels were chosen for study. The “high’‘-ferrite weld metal of composition 20Cr 1ONi 3Mo contained 16% ferrite aswelded which on heat-treatment at 850°C for 6 h transformed to sigma phase of 0.9% volume fraction leaving 1% of ferrite untransformed. After 625’C for 1000 h the sigma phase formed was reduced in amount to 0.2% with 4% of untransformed ferrite. (c) The “low-ferrite weld metal of composition 18Cr 9Ni 2Mo contained 6% ferrite as-welded, which after 85O’C for 6 h transformed to Ma& carbide with a 1.5% volume fraction and 1% retained ferrite. After 625°C for 1000 h, the carbide produced also had a 1.5% volume fraction but with 4% of retained ferrite. (d) Neutron diffraction can provide important crystallographic data on the phases present in bulk weld specimens. For example, changes in lattice parameter of both the austenite matrix and the ferrite present in the weld can be followed after heat-treatment. Thus the austenite lattice was found to contract on both the S50°C/6 h and the 625’C/lOOO h treatments. The ferrite lattice contracted appreciably more than that of the austenite, especially for the 625”C/lOOO h treatment.
References
[l] F.C. Hull, Welding J. Res. Supplement 46 (1967) 3993. [2] J.C. Borland and R.N. Younger, Brit. Welding J. 7 (1960) 22. [3] Y. Arata, F. Matsuda and S. Katayama, Trans. Japanese Welding Res. Institute 5 (1976) 35. [4] T. Boniszewski, Metals and Materials (December 1978/January 1979) 41. (51 C.A.P. Horton, P. Marshall and R.G. Thomas, CERL Note No RD/L/2045 N81: Also Int. Conf. on Mech. Behaviour and Nuclear Applications of Stainless Steel at Elevated Temps., Varese, Italy (May 1981). [6] G.J. Curtis and N. Ibrahim, Metal Sci. 15 (1981) 566. [7] B.L. Baikie, A.R. Wagg, M.J. Whittle and D. Yapp, J. Brit, Nucl. En. Sot. 15 (1976) 257.
[8] S.R. Keown 386.
and R.G. Thomas,
Metal Science,
15 (1981)
[9] F.R. Beckitt, J. Iron and Steel Inst. 207 (1969) 632. [lo] J.K. Lai and J.R. Haigh, Welding J. Res. Supplement 58 (1979) 15. [ll] G.F. Slattery and P. G’Riordan, Metallography 13 (1980) 59. [12] H. Astrom, B. Loberg, B. Bergtsson and K.E. Easterling. Metal Sci. 10 (1976) 225. [13] C.E. Lyman, P.E. Manning, D.J. Duquette and E. Hall, Scanning El. Micr. 1 (1978) 213. [14] C.E. Lyman, Welding J. Res. Supplement (July 1979) 1895. [ 151 B. Weiss and R. Stickler, Trans. Met. Sot. AIME. 3 (1972) 851. [ 161 J.E. Spruiell, J.A. Scott, C.S. Airy and R.L. Hardin. Trans. Met. Sot. AIME. 4 (1973) 1533. [17] W.E. White and I. LeMay, Metallography 5 (1972) 333. [18] J.K.L. Lai and M. Meshkat, Metal Sci. (September 1978) 415. [19] J.K.L. Lai, D.J. Chaste11 and P.E.J. Flewitt. Mat. Sci. and Eng. 49 (1981) 19. [20] L.P. Stoter, J. Mat. Sci. 16 (1981) 1039. [21] G.F. Slattery, P. O’Riordan, M.E. Lambert and S.M. Green, J. Microscopy 122 Pt. 2 (1981) 109. [22] G.F. Slattery, S.R. Keown and M.E. Lambert UKAEA Report ND-R-829(R) 1982; also to be published in Metals Technology Journal. [23] C. Windsor and C. Wright, New Scientist, Vol. 88, No. 1231 (Dee) 714. 1241 S.A. David, Welding J. Res. Supplement, April (1981) 63s. [25] N. Suutala, T. Takalo and T. Moisio, Trans. Met. Sot. AIME 10A (1979) 512. [26] A.H. Baston and D.H.C. Harris, eds., AERE Rep. No. 9278 (1978). [27] M. Sakata, R. Sinclair and C. Windsor, unpublished work ( 1980). [28] G.E. Bacon, Neutron ford, 1975) p. 96.
Diffraction
(Clarendon
Press, Ox-
[29] M.J. Cooper, K.D. Rouse and R. Sayers, unpublished work 1977. [30] M.J. Cooper and K.D. Rouse, J. Appl. Cryst. 14 (1981) 72. [31] R.G. Thomas and S.R. Keown, CEGB Report RD/M/R312. [32] C.J. Bechtoldt and A.C. Vacher, J. Res. Nat. Bur. Stand. 58 (1957) 7. [33] R.G. Thomas and D. Yapp, CEGB Report R/M/R273. [34] I. LeMay, B.J. Bassett and W.E. White, Paper Cl78 in: Creep and Fatigue in Elevated Temp. Applications (Inst. Mech. Engrs., London, 1974). [35] J.R. Haigh, Welding J. Res. supplement, 56(5) (1977) 149s.